Ion beam induced damage formation in binary and ternary
111-V compounds - an overview
E. Wendler
Friedrich-Schiller-UniversitatJena, Institut f i r Festkorperphysik, Max- Wien-Platz 1, 07743 Jena, Germany
Abstract. The damage formation in ion implanted GaAs, InP, InAs, Gap, AlAs, Al,Gal.,As and GaN is reviewed. The effect of substrate temperature, ion mass, ion dose rate and ion energy will be discussed.
In the case of GaAs, InP, InAs and GaP a phenomenological description of the observed dependences can be given basing on the
model of critical temperatures T,. The dependence of T, on the energy deposited in nuclear processes per ion and unit depth and on
the dose rate can be represented by an empirical formula. Thus it becomes possible to predict the damage evolution which is to be
expected for certain implantation conditions. The behaviour of AlAs and GaN during ion irradiation strongly differs from that of the
materials mentioned above. Only at very low ion fluences the damage formation can be correlated to the nuclear energy deposition,
whereas amorphization is achieved due to the presence of the implanted ions themselves. In A1,Gal.,As a strong dependence on the
A1 content x is found.
1. INTRODUCTION
that at that ion fluence the layers are not yet fully amorphised.
The computer code TRIM (version 87 or 97) [4] was
applied to calculate the number of foreign atoms, N,,,*,
and of primary displacements, Ndspl*,both per ion and
unit depth versus depth z. As a common quantity the
number of displacements per lattice atom, ndpa, is calculated according to ndpa = Ndspl*NI INo. NI is the ion fluence and No the atomic density of the corresponding tar* .
get. If not state otherwise, Ndspl is taken in the maximum of the distribution.
In the present paper the damage formation in ion implanted 111-V compounds is summarised. T h s topic dates
back to the early 70ies and since that many papers have
been published. T h s overview primarily emphasises the
research whch has been done at our institute. Further
references are given in the cited papers.
The paper is arranged according to the materials. The
effects occurring in GaAs, Id', InAs and GaP are discussed together in Sect. 2, because they show large similarities. In contrast to that, the behaviour of AlAs and
GaN during ion irradiation is rather outstanding and,
therefore, these materials require a separate discussion,
given in Sects 3 and 5, respectively. Section 4 is devoted
to A1,Gal.,As.
Most of our investigations were done by Rutherford
backscattering spectrometry in combination with channelling techniques (RBS). From the energy spectra of
backscattered ions the relative concentration of displaced
lattice atoms n d a versus depth z was calculated with the
computer code DICADA [l] whch is based on the discontinuous model of dechannelling [2]. A random distribution of the displaced lattice atoms withn the lattice
cell was assumed. For large defect concentrations, i.e.
when clear damage peaks occur in the spectra, the
DICADA code yields the same results as the two-beam
approximation of Bcrgh [3]. The advantages of the
DICADA code become more prominent when low defect
concentrations have to be analysed. In the following text
the relative concentration of displaced lattice atoms n d a is
referred to as defect concentration.
The ion fluences necessary for amorphisation (for short
called amorphsation-fluences) N:" were defined from
the results of the RBS measurements. At the ion fluence
NI = NI" the corresponding RBS spectrum measured in
aligned direction reaches the random level, whch corresponds to the maximum defect concentration of n d a = 1.
A comparison with other methods like transmission eleciron microscopy (TEM) or optical spectroscopy shows
2. GaAs, InP, InAs and GaP
In these materials the damage formation can be phenomenologically described on the base of critical temperatures T,. The amorphisation-fluence of various semiconductors increases exponentially with the target temperature during the implantation, TI. This dependence is
well represented by models assuming a defect annealing
within single collision cascades due to a thermally enhanced defect diffusion [5, 61. In [5] it is assumed that
each ion produces a cylindrical damage volume, the radius of which decreases with increasing implantation
temperature due to diffusion-enhanced annealing in the
outer parts of the cylinder (defect-out-diffusion model).
At the critical temperature TI = T, the amorphisationfluence reaches infinity. T, was found to depend on
Ndspl*,i.e. the density of single collision cascades, and on
the dose rate j during the implantation according to
Equation (1) was originally derived to express the dependences on Ndspl*(given in displacementsl(ion A)) and
j (given in cm-'s-') for the reversal temperature TR,
which
separates between ion-beam induced amorphsation and
cry stallisation at an amorphouslcrystalline interface [7].
CP680, Application of Accelerators in Research and Industry: 17th Int'l. Conference, edited by J. L. Duggan and I. L. Morgan
© 2003 American Institute of Physics 0-7354-0149-7/03/$20.00
670
1.0 0.8
h
i
-* \
S 0.6 C
40 keV Nt
40keVP
0 40 keV As+
0
0
- 1
&
eG
j.
-
':,
i
.B--..... ....--.+a
.......
0.4-
20 keV He+
90keVN'
v 280 keV N+
+ 200 keV Art
x 400keVGe'
A 600 keV Set
3 MeV Se+
-
f
....................
-
0..
z
GaAs, TI< 110K
0.2 -
produced by single ions. Contrary, at TI = T, the primary
produced clusters dissociate, defects recombine and anneal, leaving behnd a defect concentration of about 0.15
only over a large range of ion fluences (see below). For
TI < T, the amorphsation-fluence NI" increases with rising ion mass. After transferring NI" into the corresponding number of displacements per lattice atom, ndpaarn,
nearly constant values of 0.4 dpa are obtained for GaAs
and InAs and 0.2 dpa for InP and GaP (Fig. 1 and [S]).
Only for light ions producing more dilute collision cascades, represented by Ndspl*5 2 displacements/(ion A),
ndpa"
increases remarkably (Fig. 1). This is true also if
implantation and measurement are performed at 15 K
[S], giving a clear evidence for a-thermal defect recombination, the probability of which increases with decreasing density of the collision cascades.
-
(4
0.0 ~ " " " ' ~ " ' ~ " ' ~ " ' ~ " '
T
+
h
*
0.8
0
a
100keVB'
300 keV Sit
200 keV Art
600 keV Se+
1 6MeVArt
0
Table 1. Ranges of T, and parameJers A and B of Eq. (1) to
represent the dependences Tc(Nhspl,j ) for the different 111-V
compounds; data are taken from different authors [8].
40keVBet
v 40 keV Nt
0 40 keV Sit
A 40keVSe+
X 40keVB1+
33OkeVOt
material
GaAs
InP
InAs
GaP
Id', T,=300K
0 . 0 ' ~ " " " " " " " " " " " '
0
2
4
6
8
1
(ion A)-l
0
1
A
12651
9640
9060"
B
70.0
52.5
59.1"
range of T,
(300...360) K
(340...440) K
(300...3 15) K
(330...420) K
" data from TR(Nhsp;,j ) [30]
I
2
If the implantation temperature TI is well below T, for
the corresponding implantation conditions, the shapes of
the measured defect distributions are in good agreement
with that of Ndspl*(z)as calculated by TRIM. Of course
there is no abrupt transition from the behaviour at TI <<
T, to that at TI T,. The closer TI is to T,, the more deviate the defect profiles from Ndspl*(z)[ 131. The maximum
of the measured distributions is closer to the surface and
the final amorphous layers are thinner. The reason for
this behaviour is the defect annealing at the interface
[ 141 and the out-diffusion of defects into the
Figure 1. Number of displacements per lattice atom, ndpam,
necessary for amorphisation, versus the number of displacements per ion and unit depth, Nhsp;, for different implants into
GaAs (a) and InP (b). The data are taken from different authors, see [8].
A collection of data for T, is given in [S] and a closer
comparison of T, and T, is given in [9] for InP and InAs.
From the experimental results available so far, the parameters A and B in Eq. (1) can be determined. They are
given in Table 1 together with the corresponding ranges
of T,. In the case of the most extensively studied materials GaAs and InP, the lower end of T, correlates with
well known annealing stages: Defect annealing in GaAs
between 250 K and 350 Kcan be attributed to the mobility of Ga interstitials [lo]; in InP occurs an annealing
stage between 330 K and 420 K attributed to the mobility of P interstitials [ l l ] . This correlation between annealing stages and T, is in agreement with the defect-outdiffusion models, where the critical temperatures come
from (see above). The mobility of intrinsic defects in
GaAs, InP and GaP in the corresponding temperature
ranges was also concluded from in situ emission channelling experiments [12]. For GaAs T, is close to room
temperature. T h s explains the strong sensitivity of t h s
material to slight changes of the implantation conditions
(dose rate, temperature), which has consequences for the
reproducibility of the defect concentration remaining after room temperature implantation.
If the implantation is performed at temperatures TI <
T,, amorphisation is achieved by accumulation and overlapping of the heavily damage or amorphous clusters
0.8
0.6
GaAs: Br T,=295K
o 0.6 MeV
+ 9.0 MeV
84 0.4
0.0
&
1.0
2.0
3.0
4.0
5.0
1.0
0
-
.b
+ 0.8
8 0.6
0.4
0.0
+ 9.0MeV
8
0.2
"
m
'
m
'
L
'
b
m
0.0 0.2 0.4 0.6 0.8 1.0
displacements per atom ndpa(dpa)
Figure 2. Defect concentration in the maximum of the distribution,
versus the number of displacements per lattice
atom, ndp,, for 0.6 and 9 MeV Br implantation into GaAs at TI
= 295 K (a) and at 100 K (b) [16].
671
fluence of the dose rate can be also described withn the
model of critical temperatures. As predicted in [6] and as
follows from the empirical dependence given in Eq. (l),
the critical temperature increases with j . For a given
temperature TI an increase of j yields a higher T, and
consequently, a relative decrease of the implantation
temperature, whch results in a higher damage level.
underlying substrate, whch is proved by depth dependent measurements [ 13, 151.
For a given ion species the number of primary disdecreases when increasing the ion enplacements Ndspl*
ergy. This is caused by the enhanced straggling of the
ion range and does not effect the density of the single
collision cascades at the end of the ion path. Therefore,
the reduced defect concentration for higher beam energies at temperatures below but close to T, (Fig. 2) can be
explained by local dose rate effects [ 161.
3. AlAs
AlAs was found to be very resistant against ion-beam
induced damage formation (see e.g. [21]). Even for implantation at TI 20 K and immediate RBS measurement
at T, = TI without temperature change of the sample, the
defect concentration saturates at only ndarnaX 0.15 and
remains constant over a wide range of ion fluences (Fig.
4). Only at large fluences the amorphsation of AlAs occurs, but withn a narrow fluence range (Fig. 4).
1'
~
.s
'
'
"""'
'
'
"'""
'
'
"""II
AAs:Ar(200keV)
TI = T, = 18 K
E
ioI2
1013
1014
1015
loT7
4(cm-2)
ion fluenec N, (cm" )
Figure 3. Normalised difference in minimum yield Axmm"(representing the defect concentration) versus the ion fluence NI for
different implantations at TI = T, [8].
Figure 4. Defect concentration in the maximum of the distribution, nhma, versus the ion fluence NI for 200 keV Ar implantation into AlAs at TI = TM= 20K [24]. The different symbols
show the reproducibility of the implantations.
In the framework of the out-diffusion model, implantation at TI T, means that the reduction of the radius of
the primary damage cylinder is in the order of the radius
itself and, consequently, at least no heavily damaged or
amorphous material should remain. And indeed, this behaviour is experimentally proved. Figure 3 shows that
over a wide range of ion fluences only point defects and
point defect complexes remain, i.e. the predicted defect
annealing occurs but is not complete. Then, for very
' ~ the implanted layers belarge fluences NI > 1 ~ 1 0cm-'
come amorphous, which is confirmed by TEM investigations for InP [ 131 and GaAs [ 171. Before amorphsation a
defect band at the depth of maximum ion concentration
is detected by TEM. We believe that this defect band
suppresses the diffusion of defects into the substrate,
whch results in a defect accumulation in depths of
maximum nuclear energy deposition. Then amorphous
seeds nucleate and grow rapidly during further implantation to form a closed amorphous layer in the depth of
maximum nuclear energy deposition [ 131
The defect production strongly increases with increasing dose rate j during the implantation. The integral defect concentration remaining after implantation was
found to be proportional to j" [ 18-20]. The exponent n
and consequently the strength of the dose rate dependence increases with the target temperature during the implantation. Values of n are between 0.15 and 0.4. The in-
Although the curve nhmaX(NI)
for AlAs looks very similar to those in Fig.3, the physics behind is different. The
results in Fig. 3 were obtained for implantation at TI T,
and can be explained by a thermally enhanced defect diffusion. In the case of AlAs at TI 20 K no defects were
found to be mobile [22]. Therefore, the resistance against
ion-beam induced damage formation is an inherent property of AlAs and not caused by any thermal effects. In
[21] it was shown that the AlAs did not amorphise because the strain did not reach the critical value for amorphisation. In these studies the layer thckness was well
below the projected range of the implanted ions. In our
investigations amorphisation of AlAs (TI 20 K) occurs,
but in the depth of maximum ion concentration and not
in that of maximum energy deposition into nuclear processes [23]. The values of NI" obtained for different ion
species reveal also no correlation with the nuclear energy
density [23, 241. On the contrary, a correlation with the
atomic volume of the implanted ions was found. This is
shown in Fig. 5 which plots the ion concentration
N,,,"= NlOn*N1',
necessary to reach a defect concentration
of ndamaX= 0.6, versus the volume of the ions [24].Ths
value of ndarnaX was chosen, because the corresponding
ion fluence N," can be more precisely determined than
N,"". Our results suggest that amorphous seeds nucleate,
672
when the total volume introduced by the ions exceeds a
critical value (Fig. 5). These amorphous nuclei grow rapidly during further implantation due to ion-beam induced
amorphisation, thus causing the steep increase of ndarnaX
to nhrnaX
= 1 (Fig. 4).
b
5. GaN
Similar to AlAs, the behaviour of GaN during ion implantation is different to that for the other 111-V compounds (Sect. 2) (see e.g. [27, 281). Besides the high ion
fluences necessary for amorphisation, a pronounced surface damage occurs. To study the effects occurring in the
region of nuclear energy deposition in more detail, implantation and RBS analysis were formed at TI = 15 K
with the sample temperatures held constant [29].
Figure 7 shows the fluences dependences for different
ion species, obtained at 15 K. In Fig. 8 the depth distribution of the defect density nh(z) is plotted for different
Ar ion fluences and compared to Ndspl*(z) and Nl,*(z) as
calculated with TRIM97.
AlAs
<a
<:
Rb
\
Xe
.g
TI = TM= (15...20) K
0
0
0
1
~
1
~
10
20
30
40
1
~
atomic volumn V,, (A3)
1
number of primary displacements ndpa(dpa)
0.01
0.1
1
10
1"""'
'
""""
'
""""
'
""""
'
""'I
Figure 5. Ion concentration N,,, necessary to reach a defect
concentration of ndamax= 0.6, versus the volume of the implanted ions for implantation into AlAs at TI = TM= 20K [24].
4. A1,Gal.,As
The damage built-up in A1,Gal.,As strongly depends on
the A1 content x (see e.g. [25, 261). This is demonstrated
in Fig. 6, which shows n d p z necessary to reach
ndarnaX= 0.6 for different ion species [26]. A similar curve
for Si implantation into A1,Gal.,As is given in [25].
1
9
"
"
TI = T
'
=
"
"
j o.ol\
"
(15 ...20) K
02
04
+150 keV 0
+
, 300
, ,,,,,,,
keV Ar,
, ,,,,,,,
1014
10"
~r ion fluence N~ (cm"
,I
10l6
Figure 7. Defect concentration in the maximum of the distribution, ndamax,versus the Ar ion fluence, NI (bottom scale), and
versus the number of primary displacements per target atom,
ndp, (top scale) for ion implanted GaN. The top scale is valid
for the 150 keV 0, the 800 keV Xe and the 300 keV Ar implantations, whereas the bottom scale is valid for the Ar implantation only [29].
From our results three steps of damage formation in GaN
can be identified [29]. At the very early first stage of
damage formation, the processes are dominated by the
nuclear energy deposition (see the good agreement between the shape of nda(Z) and Ndspl*(z) in Fig. Sa). Owing
to huge recombination effects within the primary collision cascades, the defect concentration saturates at the
very low level of about ndarnaX= 0.05 (Fig. 7). We believe
that t h s value would not change, if the implantation
were not connected with the introduction of the ions
themselves. These ions seem to inhibit the recombination
of part of the primary displaced atoms by attracting and
stabilising them to form point defect clusters. This process is connected with a slight increase of the defect concentration to ndarnaX = 0.08 (Fig. 7) and a broadening of
the layers towards depth (Fig. Sa). At a certain point,
probably when the defect concentration exceeds a critical value (whch starts to reduce the recombination efficiency because of the imperfect structure of the surrounding lattice), the primary displacements again become efficient in damage production, thus causing a second stage in the damage formation. Again, the implanted
ions extend this effect, which started at the depth of
maximum nuclear energy deposition, to larger depths
00
00
,
+800 keV Xe
1013
200 keV Ar
A 400 keV Ge
2
,,,,,,,
( A
06
08
10
A1 content x
Figure 6. Number of displacements per lattice atom, ndp;, necessary to reach a defect concentration of
= 0.6, versus the
A1 content x for different implants into A1,Gal.,As [26].
For an A1 content of x 5 0 . 9 the behaviour of A1,Gal.,As
is in principle similar to that in GaAs and the shape of
the measured defect distributions is in reasonable agreement with that calculated by TRIM. For x 2 0.96 the ion
fluences necessary for defect formation increase drastically: In the case of 200 keV Ar implantation at TI = T M
= 20 K, n d p z = 18 dpa is obtained for x = 0.96 and
ndp:
= 40 dpa for x = 1 [26]. These values are much
higher than those given in Fig. 6 for x 50.9. This shows
that about 5 at.% of Ga are sufficient to brake the high
resistivity of AlAs against ion-beam induced defect formation.
673
0
estimated. In contrast to these materials, the behaviour of
AlAs and GaN during ion implantation is very different,
being dominated by huge recombination processes even
at very low temperatures and strong influences of the
implanted ions themselves. The physical reason for that
is not clear.
a
3.~10~~
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2.
03
8
.s
c
02
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01
6
8
00
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5.
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7.
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9.
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14.
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