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Analysis of Large Impurity Atmospheres at Dislocations and Associated Point Defect Reactions in Differently
n-Doped GaAs Crystals
C. Frigeri, J. Weyher, J. Jiménez, P. Martín
To cite this version:
C. Frigeri, J. Weyher, J. Jiménez, P. Martín. Analysis of Large Impurity Atmospheres at Dislocations and Associated Point Defect Reactions in Differently n-Doped GaAs Crystals. Journal de Physique III, EDP Sciences, 1997, 7 (12), pp.2339-2360. �10.1051/jp3:1997263�. �jpa-00249723�
Analysis of Large Impurity Atmospheres at Dislocations and Associated Point Defect Reactions in Dilllerently n-Doped
GaAs Crystals
C. Fkigeri (~,*), J.L. Weyher (~), J. Jim6nez (~) and P. Martin (~) (~) CNR-MASPEC Institute, via Chiavari 18/A, 43100 Parma, Italy
(~) Fraunhofer-IAF, Tullastrasse 72, 79108 Freiburg, Germany
(~) Universidad de Valladolid, Departamento de Fisica de la Materia Condensada, ETS Ingenieros Industriales, 47011 Valladohd, Spain
(Received 2 December 1996, accepted 21 August 1997)
PACS 61 72.Ff Direct observation of dislocations and other defects (etch pits, decoration, electron microscopy, x-ray topography, etc.
PACS.61.72 Yx Interaction between different crystal defects; gettering effect PACS.71.55 Eq III-V semiconductors
Abstract. The large impurity atmospheres at dislocations typical of n-type (Si- or Te- doped) GaAs crystals have been analysed by localized measurements of the free electron con-
centration, diffusion length and DSL etching velocity. The atmospheres always contain dopant
atoms as well as point defects (complexes) whose formation and type depend on the type of
dopant impurity and melt stoichiometry. The donor- or acceptor-like characteristics of such point defects (complexes) are responsible for the observed remarkable difference m the electrical and recombinative properties of the atmospheres between the differently doped crystals. The point defect reactions at the base of the formation of the slip traces are discussed The possible mechanisms of the impurity-dislocation interaction leading to the formation of the atmospheres
are also considered
1. Introduction
The interaction between dislocations and impurities causes the formation of atmospheres
around dislocations and is thus one of the main sources of non-uniform distribution of im- purities in semiconductors which produces inhomogeneities of the electrical and optical prop-
erties. Impurities also affect dislocation generation and the overall dynamic properties of the dislocations, see reference ill and references therein.
As shown below, in as-grown bulk n-type GaAs the impurity atmospheres at dislocations
are often very large (> 5 ~m) and have different electrical and recombinative characteristics, depending on the dopant species and melt stoichiometry of the sample. In order to understand
the reasons for these differences, this paper reports on the analysis of such large atmospheres in
GaAs doped with different n-type impurities (Si and Te) and grown under different stoichiom- etry deviations. Analyses have been carried out by measurements of relevant GaAs properties,
(*) Author for correspondence (e-mail: Frigeritlmaspec.bo.cnr.it)
@ Les (ditions de Physique 1997
like majority carrier density, photoetching velocity and minority carrier diffusion length, in the large atmospheres. The high resolution (~u I ~m) techniques used for the analyses, i.e.
photoetching, SEM-EBIC (scanning electron microscope-electron beam induced current) and microRaman, are briefly reviewed.
2. Experimental
The (100) GaAs samples were either Si-doped (n
= 1-15 x 10~~ cm~~) or Te-doped
(n'= 2 x 10~~ cm~~) Two types of Si-doped crystals have been investigated, i.e. either grown from As-rich melt or from Ga-rich melt (XAS " 0.494). To asses the doping level concentration and the recombinative properties at the large impurity atmospheres associated with the dislocations, three experimental techniques were used, namely DSL photo-etching, SEM-EBIC and microRaman spectroscopy. Previously published data [2, 3j on high spatial
resolution photoluminescence (PL) mapping and some local spectral measurements have also been introduced in the discussion of the results. TEM was also employed. Te-doped speci-
mens for TEM were first photo-etched to assure correlation between surface features and the
underlying defects.
a) DSL etching.- DSL etching is based on a Diluted Sirtl-like (DS) solution (Cr03.HF:H20) [4-7j. The electroless etching of GaAs in the DS solution occurs by a mechanism of oxidative
dissolution of the GaAs molecule that requires six holes (h+) per unit molecule [5, 6j:
GaAs + 6h+ ~ Ga~+ + As~+
The holes are supplied by the etching solution [5,6j. The etching rate is, however, greatly
increased by illuminating the sample with Light (DSL etching) [4-6j. The dissolution of the GaAs molecules is then mostly due to the holes generated by light inside the sample that are able to reach the surface [6-8j. Their density, ps, depends on the band bending at the interface
between GaAs and etching solution (Fig. I) according to [8]
p~ = loo + pscR + Pb) exPlqKc/kT)
where po is the equilibrium bulk hole density, pscR the density of the holes photogenerated
inside the surface space charge region (SCR), pb the fraction of the holes photogenerated in the neutral bulk that reach the surface SCR, Kc the surface band bending potential, q the electron charge, k the Boltzmann's constant and T the temperature. On the other hand, the density of the equilibrium and photogenerated electrons at the surface is reduced by exp(-qfc/kT), so
that the holes at the surface suffer very little recombination with electrons Any reduction of holes at the surface reduces the etch rate which results m the formation of morphological reliefs, such as hillocks, on the etched surface This occurs, for instance, at extended or point defects which recombine holes [4,7,8j or when the width of the surface SCR decreases, i-e- the doping level increases [8,9), (Figs. 1b, c). Of course, the opposite is also true, i-e- additional supply of holes from sources inside the sample increases the etching velocity and surface depressions
are formed. The DSL etching rate is therefore extremely sensitive'to the (changes of) doping concentration and to the recombinative properties of the defects[ As regards defects, DSL
works like the techniques based on the injection and detection o( minority carriers, such as
EBIC. The resolution of the method is in principle atomic, but in' practice it depends on the resolution of the technique used to evaluate the surface after etching. For optical microscopy used in the differential interference contrast (DIC) mode the reso1~ltion is limited to
r~J 1 ~m.
The height profile of the etch features was determined by a step profilometer and by phase
/ n~
~~ ~
n
.$~~~ i~2~ ~~
DSL etched surface or;g;nal surface
b) C)
Fig. i a) Sketch of the DSL etching mechanism of GaAs b) A recombmative defect, like a
dislocation (I), decreases the etch rate and gives rise to a hillock. c) Differences of free electron
(doping) concentration produce differences m the etch rate (depletion region effect), the n regions
being etched faster than the n+ regions See text.
stepping microscopy (PSM) in association with a computer for the 3D reconstruction of the etch features from phase changes [10).
b) EBIC.- Measurements were performed by the energy-dependent method with the Schot-
tky junction planar geometry. In such method the EBIC collection efficiency is a function of both diffusion length L and width W of the depletion region associated with the Schottky diode, as well as of the beam energy [11,12]. EBIC efficiency increases when either L or W, or
both, increases. The values of W and L are extracted by best fitting the experimental efficiency
versus beam energy curves to the theoretical curves [11,12). From the experimental value of
W the free electron density n
= ND NA, with ND and NA the ionized donor and acceptor
density, respectively, is obtained by the formula [12j
n = 2eVo/qW~ (1)
where e is the GaAs dielectric constant and % the built-in potential of the Schottky diode.
On the other hand, L is related to the deep traps by the relation
L
= (D/vthatNt)~/~
with D diffusion coefficient of the minority carriers having a thermal velocity vth, at and Nt the capture cross section and density of the deep trap, respectively. The value of L m GaAs is, in fact, much lower (<
r~~ 1-2 ~m) than that predicted by the band-to-band recombination
theory for all doping levels, so that the dependence of L on the shallow doping density can be
neglected [13). The spatial resolution of the method is
r~J 3-4 ~m The beam current used was
< 1 nA and the beam energy was varied between 5 and 40 kev.
c) MicroRaman.- The backscattering first order Raman spectrum of a randomly oriented
undoped GaAs sample exhibits two phonon peaks, I e. the longitudinal optical (LO) and the transverse optical (TO) peaks. Due to the Raman scattering selection rules [14j, m an exactly
oriented (100) GaAs crystal only the LO mode is allowed. The presence of the TO mode is indicative of crystal misorientation m the sampled area with respect to the (100) orientation
or of the presence of defects that activate the TO mode. MThen the doping level n exceeds
r~J 3-4 x 10~~ cm~3 the free electron gas strongly interacts with the optical phonons through
their macroscopic electric fields causing a splitting of the LO mode into two additional branches, L- and L+ having "requencies UJL- and UJL~, respectively, with values 0 < UJL- < uJLo and uJLo < UJL+ < uJp, where uJLo, uJp are the LO and plasma frequency, respectively [15). The LO band arises from the surface depletion layer. The ratio between the intensity I of the LO and L- Raman peaks is related to the width d of the surface depletion layer by [16]
IILO)/IIL-)
= Ale~"~ i) 12)
where A
= 2.65 for GaAs and a is the absorption coefficient for the excitation light. The free electron concentration is still calculated by equation (1), by using for % the value of the
GaAs lair system and by replacing W by d.
For the microRaman experiments the Ar+ laser light source emitting at ~
= 514.5 nm was
employed. By focusing the light beam onto the sample surface by an optical microscope a
spatial resolution of <r~J 1 ~m is achieved.
3. Results
The dislocations, and related large (r~J 5 -50 ~m) impurity atmospheres, examined in this paper
are so-called grown-in (G) and grown-in, then moved by stress (G-S) dislocations, that were discussed in detail elsewhere [2,7,17,18]. G-S dislocations are those dislocations that, during cooling down, have been moved by stress from one start position S, where they were sitting as G dislocations, to another end position E, very often leaving behind a recombinative trace (T)
of impurities [2,7,17,18). In the Te-doped samples only G dislocations have been detected.
3.1. As-RICH, Si-DOPED GaAs. Figure 2 shows the DIC optical image after DSL etching
and the EBIC image of typical G dislocations in Si-doped GaAs grown from As-rich melt.
The outcrop of the dislocation is the small summit on top of the big DSL-revealed hillock, as
evidenced by the step profile (Fig. 2b). In the EBIC image such small tip corresponds to the small and darker dot in the middle of the large dark atmosphere, whose presence is clearly seen
in the EBIC line scan (Fig. 2d). At such small dot the non-radiative recombination is higher
than in the surrounding cloud.
i~_,
-D
50 nm
~' lo pm
b
&
y~
/. ~
? c I
lo pm
C D ~
. .
d
Fig. 2 G dislocations m Si-doped GaAs grown from As-rich melt. a) DIC image of a DSL-etched
sample. b) Step profile along the horizontal diameter of the hillock in a). c) EBIC image d) EBIC line scan along the horizontal diameter of the atmosphere in (c) D is the outcrop of the dislocation.
In a) bar
= 20 ~m.
Deconvolution of the EBIC line scan of Figure 2d shows that the FWHM (full width at half maximum) of the EBIC profile across the small darker dot in the middle of the large atmosphere is similar to the FWHM of the EBIC line scans taken across dislocations without large atmosphere [19, 20) and corresponds to the FWHM predicted by Donolato's theories [21) of EBIC contrast at dislocations
This agrees with the expected value of the diameter of the recombination cylinder around
dislocations, due to the dangling bonds at the dislocation core, as can be calculated from the Labusch-Schrbter [22, 23) or Read [24) theories. In the Labusch-Schroter theory, by defining
the diameter DLS of the recombination area around a dislocation as 2~, where ~ is the Debye screening length of that theory [22,23j, one has DLS
" 2~
= 2[kTe/q~ (n p)j~/~ = 11 nm, for q r~J 5 x 10~~ cm~3 » p, p being the hole density, with k, T, e, and q already defined. In
Read's theory [24j, the recombination area effective for recombination of the minority carri-
ers at a dislocation is the space charge cylinder of radius r that forms around a dislocation in order to neutralize the charged dangling bonds at the dislocation core. Its diameter is DR ~ 2r
= 2[f/7rc(ND NA))~/~, where f is the occupation fraction of the dangling bonds
with distance c between them With ND NA
" n = 5 x 10~~ cm~3, c
= 0.5 nm and f
= 0.10, it is DR
" 2r
= 23 nm.
DLS or DR give the actual size of the recombination cylinder around a dislocation. They are
. E
)
50 nm
'
Fig. 3. G-S dislocations m As-rich, Si-doped GaAs. a) DIC image of a DSL-etched sample. b)
Step profiles across S, T and E in jai D is the outcrop of the dislocation. c) EBIC image. d) PSM image of a G-S dislocation m a sample with average n = 4.5 x 10~~ cm~~ Bars
= 20 ~m.
more approximated by the size of the small darker dot rather than by the size (r~J 5-50 ~m)
of the large impurity atmosphere around it. The fact that the EBIC FWHM is greater than DLS and DR is due to the spatial resolution of the EBIC methodl that is determined by the diameter R of the electron beam-specimen interaction volume rather than by the defect size
as far as the latter is smaller than R [21j. The fact that only the small summit corresponds to the true dislocation was confirmed by TEM.
DSL etching, EBIC and microRaman have been applied to the study of the large impurity atmospheres (interaction zones) only.
In Figure 3 optical microscopy, after DSL etching, and EBIC images of G-S dislocations
10~° 0.tt i~lo ~ ~
~_
?
+- -~ L
/ ~ ~
0.6 ~ ~~
~
o 6 ~
"j / ~
C ~
~
~
i
c
~
C °.4 o,4 ~'
10~~ j
M
~ M
10~~
~
0.2 0.2
0 10 20 30 40 0 10 20 30 40
a) Distance X from defect centre (~m) ~) Distance x from defect centre (~m)
Fig. 4. Si-doped GaAs grown from As-rich melt. Plots of the free electron concentration n and diffusion length L, as measured by EBIC, as a function of the distance from the defect centre for a)
the start point and b) the end point of a G-S dislocation. C
= centre of the atmosphere, H
= halo, M
= matrix (X > r~J 20 Mm). Average n in the sample 2.2 x 10~~ cm~~.
are given. DSL etching produces hillocks at all atmospheres as well as at the trace. Such atmospheres and trace appear dark by EBIC. The step profiles show that there is no dislocation at the S point (Fig. 3b), as also seen by EBIC [17,19, 20j. EBIC images show that at the E point of the G-S dislocations the impurity atmosphere is surrounded by a halo with bright
contrast (typical dot-and-halo feature), which is indicative of higher values of L and/or W
with respect to the surroundings. The origin of the bright halo was discussed elsewhere [17j.
The bright haloes were also observed sometimes at G dislocations.
EBIC measurements of n and L show that in the atmospheres at the start and end points
of a G-S dislocation there is an increase of n and a decrease of L with respect to the matrix
(Fig. 4). Similar results have been obtained for G dislocations. On the other hand, in the
bright halo around the end point there is an increase of L and a decrease of n, that indicates
a depletion of donors.
The increase of n in the atmospheres is confirmed by microRaman (Fig. 5). The increase is higher at the end point of the G-S dislocations, as also seen by EBIC. No change of n was
detected by either microRaman or EBIC at the trace of the G-S dislocations
3.2. Ga-RICH, Si-DOPED GaAs. Figures 6, 7 show images of G and G-S dislocations, respectively, in Si-doped GaAs grown from Ga-rich melt. The areas around dislocations exhibit features different from those seen in As-rich crystals. Such areas correspond to depressions after DSL etching, as confirmed by the step profiles, whereas by EBIC they exhibit bright contrast
instead of dark. In the case of G dislocations and end point of G-S dislocations these areas
are made up of two parts, i-e-, an outer annular halo with brighter EBIC contrast (deeper
DSL depression) and an inner area with less bright EBIC contrast (less deep DSL depression) (Figs. 6, 7). The external halo is equivalent to the bright halo observed in the As-rich samples,
whereas the inner part corresponds to the impurity atmosphere, which has therefore opposite
characteristics to those in the As-rich samples (Figs. 2, 3). The S point of the G-S dislocations has no external halo, as in the As-rich samples On the other hand, the trace always yields a ridge-like etch-feature in the DSL bath, similarly to what observed in As-rich samples (Fig. 7b).
In the EBIC image (Fig. 7c) only the bright contrast at the S and E points can be seen. The
trace appeared dark by EBIC on the SEM display, as for the As-rich sample, but the contrast
was so weak that it could not be transferred to the picture. The outcrops of the dislocations
E
~
m'~ fl~
o E
O
~ ~
'O g 'O
M 5
w -
c c
~ T
0 25
50 °
a Distance (pm ) b Distance (pm)
Fig. 5. Si-doped GaAs grown from As-rich melt. Values of n calculated from MicroRaman spectra,
as a function of the distance across the defect and parallel to the trace, at a) the start point S and
(b) the end point E of a G-S dislocation. In a) the atmosphere centre is at X
= 0 ~m. T
= trace, M
= matrix. Average n in the sample 4.5 x 10~~ cm~~
produce small hillocks after DSL etching (tip D in the step profile of Fig. fib) or appear as dark dots by EBIC, as in As-rich crystals, due to the intrinsic property of the dislocation core which
always act as a non-radiative recombination cylinder, because of the dangling bonds or land of their contamination by impurities. Figures 6, 7 confirm that the large impurity atmospheres
around dislocations have characteristics that are different from those of the dislocation cores.
In the EBIC bright areas (DSL-revealed depressions) at G dislocations and end points of a G-S dislocation, there is a strong decrease of n and an increase of L with respect to the matrix, the changes being more pronounced in the halo (Fig. 8). This result agrees with the similar result obtained in the EBIC bright halo at the end point of G-S dislocations in As-rich GaAs
(Fig. 4). A decrease of n and an increase of L was also observed at the start point of G-S dislocations.
Figures 9a and b give the PL spectra, taken at 4 K with a spot size of r~J 2 ~m, in the annular halo and in the central area (atmosphere) inside the halo at the end point of a G-S dislocation [2j. They show that in the external halo there is a decrease of SiGa whereas in the central atmosphere there is a remarkable increase of SiGa and a decrease of SiAs with respect
to the matrix [2j. A similar result was obtained for G dislocations [3j and S points of G-S dislocations [2).
3.3. Te-DOPED GaAs. The Te-doped material exhibits characteristics similar to the Ga- rich crystal, I. e. depressions are created by DSL etching m the surroundings of the dislocation
outcrops (Figs. 10a, b). Inside such depressions several small etch hillocks are detected (e.g., tip D in Figs. 10a, b) that are due to the dislocation as well as to the loops surrounding it (see
TEM results below). The area surrounding the depression is depleted of etch features over a distance of
r~J 5-10 ~m. Besides this, many small hillocks are visible all over the DSL etched surface.
The areas surrounding the DSL-revealed depressions correspond to areas of EBIC bright
contrast, whereas the region of the depressions with the small hillocks give rise to EBIC
