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Point defects in III–V materials grown by molecular beam epitaxy at low temperature
16
Materials Ai'ieHce am/t:)tgitwering, B22 1993 16 22
Point defects in III-V materials grown by molecular beam epitaxy at low temperature P. Hautojfirvi, J. Miikinen, S. P a l k o a n d K. S a a r i n e n
Laboratory of Physics, Helsinki Universityof Technology, SF-02150Espoo (Finland) C. C o r b e l a n d L. L i s z k a y
Centre d'Etudes NuclOairesde Saclay, INSTN, F-91191 Gif-sur-Yvette (France)
Abstract The present understanding of the point defects in GaAs and InP grown by molecular beam epitaxy at low temperature (LT) is briefly reviewed. New results on vacancies and ion-type acceptors obtained by positron annihilation are given. Depending on the growth temperature, Ga vacancies or small vacancy clusters are seen in LT GaAs in the concentration range l0 ts cm -3. No signal from Ga antisites is found. The LT InP layers contain vacancies, identified as In vacancies, in the concentration range 10 ~ cm 3. Ion-type acceptors, probably In antisites, are also seen in concentrations of l0 w cm-3. The annealed layer contains small vacancy clusters.
1. Introduction Low temperature (LT) growth of thin III-V layers has recently attracted much attention as a potential technique to obtain semi-insulating layers [1, 2]. Such layers are strongly needed for various micro- and optoelectronic device and integrated circuit applications. The layers contain point defects in extremely high concentrations. To identify and understand their role in the electrical and optical properties of the LT layers is a new challenge for various defect spectroscopic techniques. In LT GaAs and InP the main defects seem to be the group V antisite atoms AsGa and P~n, which have been identified as native donors [3, 4]. However, there is little information on native acceptors. Vacancies and antisites of group III atoms are generally believed to be acceptors. Positron annihilation is a unique method to identify vacancies [5-7]. In addition, as positive particles, positrons are also sensitive to ion-type acceptors. Here we present new information on the native acceptor defects in LT GaAs and InP by using positron annihilation. In this paper the main properties of the point defects in LT GaAs and InP are briefly summarized in Sections 2 and 3. The positron beam spectroscopy of thin layers is described in Section 4. The results on LT GaAs and InP in Sections 5 and 6 show that the layers contain high concentrations of native vacancies (V) tentatively identified as V(~a and VI,, respectively. LT (1921-5107/93/$6.00
InP layers also contain ion-type acceptors, most probably In antisites. The paper is summarized in Section 7.
2. Point defects in LT GaAs The defect properties of LT GaAs depend strongly on the growth temperature Tg, which varies from 200 to 300 °C. Analysis by particle-induced X-ray emission (PIXE) shows an As excess in the layer depending on the growth condition. The As excess is up to 1.5 at.% at Tg=190 °C and decreases monotonously as Tg increases. Layers grown at 300 °C or higher temperatures are stoichiometric within the sensitivity limit of 0.1 at.% of the PIXIE method [2, 8]. Studies by Rutherford backscattering and channelling indicate that As interstitial complexes rather than isolated As interstitials are present [8]. The only well-identified defect, however, is the arsenic antisite ASGa detected by the electron paramagnetic resonance (EPR) technique in its positive charge state ASGa÷ [3, 9]. Its concentration is 3-5 x 1018 cm -3 at Tg= 200 °C and decreases by an order of magnitude when Z. is increased to 300 °C [10]. The neutral state ASGafi is observed by optical absorption in the near-IR region [11]. Its concentration is a maximum of 2 × 10 2o c m -3 at Tg= 190 °C and decreases to 2 x 1019 cm -3 at Tg = 270 °C [2]. An open and controversial question is the photoquenching © 1993 - ELsevier Sequoia. All rights reserved
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Point defects in L T-MBE III-V materials
property of the AsGa defect in the LT layers, which seems to be different from that of the ASGa defect in semi-insulating bulk GaAs. Kaminska and Weber [2] have offered an interesting point of view. The strain in LT GaAs layers changes the local hyperfine field. The different properties of these two defects can be ascribed to differences in the local fields seen by the two defects rather than to differences in their nature. More generally, this question addresses the relation of the ASGa antisite to the EL2 centre in GaAs. Evidently ASGa is the main native donor in LT GaAs. It is activated at concentrations well above 10 ~8 cm -3. To compensate the ASGa+ centres, native acceptors are needed. Evidence of an intrinsic acceptor level at Ev + 0.3 eV has been given [12], which could be due to Ga vacancies. Yu et al. [13] have suggested that a complex seen by photolumiscence could be a VGa-As i pair. Measurements of the localized IR vibration modes of Si atoms suggest that Si-WGapairs are formed in Si-doped layers [14]. We show later in this paper that positron annihilation gives direct evidence of Ga vacancies at concentrations sufficiently high to account for the compensation. The conductivity of the as-grown layer is due to hopping within a defect band of ASGa antisites [15, 16]. After annealing of the layers at 600 °C, which always occurs in device applications, the normal conduction mechanism through thermally excited carders appears. The layers have become semi-insulating but the carrier lifetime and mobility are still low. The major annealing effect on the microstructure of the layers is the formation of As precipitates of size 6 nm to accommodate the As excess [17]. There is also a reduction in the concentration of ASGa° to below the detection limit of 1018 cm -3. In addition, the ASGa+ concentration drops to the level of 1017 cm-3 [9]. Two models have been suggested to explain the semi-insulating conductivity after annealing at 600 °C. (1) The point defect concentration is still high enough to explain the low carder lifetime and mobility. However, the question of the concentration of ASGa° and of other defects such as VGa remains open. (2) The high resistivity is due to metallic As precipitates which act as buried Schottky junctions. The space charge regions generated by the Schottky barriers overlap in the matrix, which is thus completely depleted of free carriers [18].
3. Point defects in LT InP
The success in the case of LT GaAs layers has stimulated efforts to grow LT InP and understand its electrical and structural properties. The growth temperature ranges from 130 °C upwards. The inter-
17
esting region is between 170 and 200 °C [19]. The amount of P excess is about 1 at.% at Tg= 170 °C and decreases with increasing growth temperature. The layers exhibit highly n-type conductivity except when they are passivated by hydrogen in the as-grown state [19, 20]. Beryllium doping does not compensate layers grown at Tg<300 °C [4, 19, 20]. Layers annealed at 600 °C stay conductive, although the P excess precipitates into small particles of size 3-7 nm [21 ]. The electrical properties of LT InP layers indicate that the dominant intrinsic defect is present in concentrations of 1019 cm -3 [20]. EPR measurements are not possible owing to the high conductivity of the layers. Recently Dreszer et al. [4] have found two deep donor levels at 0.11 eV above 0.23 eV below the conduction band. Optically detected magnetic resonance (ODMR) combined with photoluminescence reveals that the two levels are the first and second ionization levels of the P antisite double-donor defect. The upper level being resonant with the conduction band gives rise to the high free-electron via auto-ionization. To explain the photoluminescence spectrum, the authors need to call for an unidentified donor level at 0.8 eV above the valence band. We will see in Section 6 that in addition to the group V antisite atoms which seem to be the prevalent defects in LT InP, positron annihilation experiments indicate the presence of vacancies associated with the In sublattice in concentrations of 10 ~8-10 ~9 cm- 3
4. Positron annihilation technique 4.1. Positron dynamics in semiconductors
Energetic positrons in solids thermalize within a few picoseconds. In semiconductors the mean lifetime of free positrons is typically 250 ps. The motion and trapping of thermal positrons are in many respects analogous to those of free carriers apart from the "heavy" positron effective mass m*+ = 1.5m e. The positron diffusion coefficient at 300 K is around 2 cm 2 s- 1 [22, 23]. The diffusion is limited mainly by acoustic phonons, leading to the T -1/2 temperature dependence. Positrons feel electric fields and have a mobility of about 100 cm 2 V-1 s-1 t 300 K with a T -3/2 dependence. The Bloch state positron wavefunction is squeezed by the positron-ion repulsion into the interstitial regions between atoms. Thus an open-volume defect such as a vacant lattice cell is an attractive centre where positrons get trapped. Positron trapping is analogous to carder capture [24]. The trapping rate r is proportional to the defect concentration, r =/~ Cd. The charge state of a vacancy has a strong effect on the trapping coefficient/~. For a neutral vacancy/~ ~- 1014-1015 s-l
1N
P. Etautojdrvi et al.
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Point defects in L T-MBE 111- V materials
independently of temperature. For a negative vacancy ~t ~ 1015-1016 S I and it increases by an order of magnitude at low temperatures. Positive vacancies do not trap positive positrons. Minimum vacancy concentrations required to trap positrons during their lifetime are about 10 l 5_ 1016 cm- 3. Owing to their different electronic environment, the annihilation characteristics of trapped positrons are different from those of free positrons in the bulk. The reduced electron density at a vacancy increases the positron lifetime by 10%-20%. Because of the smaller overlap of the trapped positrons with core electrons, the momentum distribution of the annihilating pairs becomes narrower. The change in the charge state of a defect causes measurable changes in the annihilation characteristics, thus giving the possibility to determine the ionization levels of the defect in the band gap [25]. Owing to coulombic interaction, a positron may also get trapped at weakly bound Rydberg states around negative ions. Thus information can also be obtained on ion-type acceptors [26].
4.2. Positron beam analysis Positrons from a radioactive source first hit a separate moderator crystal which has a negative work function for positrons. Thermalized positrons close to the moderator surface are spontaneously emitted into a vacuum. They form a beam which is accelerated and guided to the sample. In an experiment monoenergetic positrons are implanted at various depths in the sample by varying the beam energy E typically from () to 30 keV. In GaAs this energy range scans depths from 0 to 1.5 ~m. The positron-stopping profile is, however, rather wide [27], limiting the depth resolution of the technique. For reviews on positron beams see refs. 5 to 7. Since the start signal from the positron emission is not available in a beam, the positron lifetime cannot be measured, but the 511 keV line of the annihilation radiation is recorded. This line shows considerable Doppler broadening due to the motion of annihilating pairs. It is customary to use a line shape parameter S (or W ) defined as the ratio of the counts in the central (or wing) region of the annihilation line to the total number of counts in the peak. Owing to their high momentum, mainly core electrons contribute to the wing parameter W, whereas annihilations with the valence electrons fall predominantly in the energy region of the central parameter S. Both line shape parameters have characteristic values for each material, depending on the electron momentum distribution. When positrons are trapped at vacancytype defects, the annihilations with core electrons are reduced with respect to the perfect lattice, increasing the S parameter and decreasing the W parameter.
Hence the line shape parameters can be used to characterize vacancy defects in solids. The measured line shape parameter as a function of beam energy, S(E) (or W(E)), is a superposition of the characteristic values S~ for various annihilation states i. As an example, let us assume a defect layer in the nearsurface region of a sample. There are annihilations at the surface (surf), in the defects (def) and in the bulk (bulk) and we have
S(E) = F~urf(E) Ssurf~-/~\tcf(E) Sdef + Fbulk(E) Sbulk
(l)
Here b](E) denotes the fraction of annihilations in state i. To calculate the fractions Fi(E), the positron diffusion equation must be solved. Analysis of the experimental results is carried out by varying the positron-trapping profile ~c(x)=/zC~(x) in the diffusion equation until a satisfactory fit between the experimental and calculated S(E) is reached. The line shape parameters S and W depend on the fractions of positrons in various states. To overcome this difficulty, a new defect-specific parameter R is defined as [28, 29] R=
IS-&u,kl I W- Wbu,kI
(2)
If there are only two annihilation states, e.g. the bulk and a defect, R is independent of the fractions F~(E) and its value can be used to identify the defect.
5. Results on LT M B E GaAs
5.1. Samples The GaAs layers were grown in a Varian 360 system under normal, As-stabilized conditions at a growth rate of 0.8 /zm h-~ on semi-insulating GaAs substrates. A layer 4 /zm thick was grown at 200 °C. Two other layers of thickness 2.5 p m were grown at 225 and 300 °C and doped with Be to around 8 x 1019 cm -3. One sample was furnace annealed for 30 min in vacuum under face-to-face protection at successive temperatures of 300, 400 and 500 °C.
5.2. Results Figure 1 shows the experimental Doppler-broadening parameter S as a function of the positron beam energy E in three layers grown at 200, 225 and 300 °C. As reference sample we use a piece of semi-insulating LEC-grown bulk GaAs wafer which is known to contain no positron traps. The change in S(E) for E < 5 keV represents the transition from surface annihilations to annihilations in the layer. The flat part above
P. Hauto#irvi et al.
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Point defects in L T-MBE III- V materials
MEAN STOPPING DEPTH (pro) 0.1 0.5 1.0 1.5 I
I
I
I
I
m
¢n 1.04
I
1.04
-o-¢
,,--
-
,--o--*•
1.03
=,,, = l =
225"C
o o-_.
~1.02
.
..
o_,_o
io.1 . 0 ~ LU
,t
#_o~~-
•
/
"P.
I
~ 1.01:Or)
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To = 200"C
cc uJ
TG = 200"C
19
j300*C
I'~
,.ooll
--1.00
300"C
,
,
,
i
o 0
10 20 INCIDENT ENERGY (keV)
i
I 200
GaAs ref. i I 300
(K)
Fig. 2. Temperature dependence of S parameter measured at 15 keV positron energy in LT GaAs grown at various temperatures Tg.
cm-3. Above 5 keV implantation energy the positron trapping is saturated and the S = 1.037 value is characteristic of the defects in this layer. The defects are divacancies or small vacancy clusters [28, 29]. We performed furnace annealing of this layer up to 500 °C. No changes in the S ( E ) curves except at the surface (E < 2 keV) could be seen. This shows that the vacancy clusters are stable and difficult to anneal out. The high stability of vacancy clusters has also been observed earlier in ion-irradiated GaAs [28]. It is important to note that the total thickness of this layer is 4 #m, whereas positrons probe only the uppermost 1.5 # m region. Thus it is still open whether the region within the critical thickness (about 2 #m) from the substrate also contains vacancy clusters as the dominant defects. Figure 2 shows the temperature dependence of the S parameters in the layers measured at 10 keV positron energy. The message is clear: no temperature dependence. This means that the vacancy signal is not reduced at low temperature by positron trapping at Rydberg states around negative ions. This is in contrast to what has been observed in electron-irradiated GaAs, where in addition to Ga vacancies, negative-iontype acceptors identified as Ga antisites have also been observed [26]. We can conclude that the concentration of ion-type acceptors must be at least an order of magnitude less than that of vacancies. To determine whether the vacancies in the Tg-225 °C layer are in the Ga or As sublattice, we refer to their charge state. In semi-insulating GaAs, Ga vacancies are acceptors with a 2 - or 3 - charge, whereas As vacancies are positively charged [25]. Since positrons do not get trapped at positive defects, the traps can only be Ga vacancies. Such an identification a n d 1019
5 keV is characteristic of layer annihilations. To facilitate comparisons, the level of the S parameter in the reference sample has been normalized to 1.00. We can see a perfect overlap at the S parameter levels in the Tg=300 °C layer and the reference sample. This means that no vacancies in the layer are seen. Also, the transition from the surface value to the layer value extends up to 10 keV, indicating freepositron diffusion. The Tg= 225 °C has an S value of 1.023. This is well above the reference bulk level, which indicates the presence of vacancy-type defects in the layer. The very rapid increase from the surface S value to the plateau means that the positron diffusion length L + ~ 100/k is very short compared with the free-positron diffusion length L+ ~ 1600 A [23]. This indicates a high concentration (about 1019 cm-3) of defects trapping positrons, which means that all positrons between 3 and 20 keV annihilate when trapped. The S value of 1.023 is characteristic of the defects in the layer. This value is close to those found earlier for As and Ga monovacancies. Thus we can conclude that the dominant traps are monovacancies in this layer. The decrease in S above 20 keV is due to the increasing fraction of positrons, which at higher energies penetrate into the substrate with a negligible defect concentration compared with the overlayer. The S value in the Tg= 200 °C layer is 1.040. This is even higher than the value for monovacancies. The increase in S between 0 and 5 keV gives an effective positron diffusion length L+ = 1 7 0 /k. The corresponding positron trap concentration is between 10 ]8
I I00
TEMPERATURE
30
Fig. 1. Line shape parameter S as a function of positron beam energy in LT GaAs layers grown at various temperatures T~. The Tg--300 °C curve coincides with that of the semi-insulating GaAs crystal (not shown) used as reference.
o.~o__~:~o-g-
il,~-,:W
20
P. Hautojdrvi et al.
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Point defects in L T-MBL 111- V materials
leads us to conclude that the vacancies act as native acceptors. It is interesting that in the layer grown at 300 °C no Ga vacancies are seen. This means that the concentration of Ga vacancies is less than 101~' cm -3. Von Bardeleben et al. have measured the concentration of ASGa+ as a function of growth temperature. Their results show the same trend as for our VG.: when Tg is 300°C or higher, the AsG~+ concentration has decreased by more than one order of magnitude from the layers grown at 200 °C [10]. In conclusion, positron annihilation shows the existence of vacancies and vacancy clusters in concentrations of 1018-1019 cm -3 in the Be-doped layer grown at 225 °C and in the undoped layer grown at 200 °C respectively. No signal from Ga antisites is observed. The vacancies are associated with V~a and so they can be the acceptors compensating the native Asoa donors observed earlier by EPR. In the Be-doped layer grown at 300 °C the V ~ concentration is below 1016 cm-3.
M E A N S T O P P I N G DEPTH (/am) 0.1 0.5 1.0 1.5 I
I
l
I
~oO°%°°°°°-~°Oo~o
m 1.04 nuJ
o
~ 1.03 <
°°- TA -- 5600C O
~ 1.02 w •
"I" 1 . 0 1
~
AS-GROWN
w Z
~ 1.00
InP ref. IA
0
I
I
10
20
=
I
30
INCIDENT E N E R G Y (keV)
Fig. 3. Line shape parameter S as a function of positron beam
energy in LT InP layers in as-grown state and after annealingat 560 °C. As reference sample an [nP (Zn:2 × 10~s cm -3) crystal has been used.
6. Results on LT MBE InP
6.1. Samples Several InP layers were grown by gas source molecular beam epitaxy at Thomson, Corbeville. For details of the growth system see ref. 19. The growth temperature was 180 °C and the thickness of the layers 1.5/~m. The PH 3 flow rate was 8 or 22 sccm (standard cm 3 min- 1). Most of the layers were doped with Be to 5 x 1017 cm-3. All the layers were n type with a carrier concentration of 5 x l0 t6 cm -3 even after annealing at 560 °C for 30 min. As a "defect-free" reference we used a Czochralski-grown InP (Zn: 2 x 10 Is cm-3) bulk wafer which, according to positron lifetime measurements, does not exhibit any positron trapping by defects.
6.2. Results Figure 3 shows the experimental Doppler-broadening parameter S as a function of the positron beam energy E for three different samples. The InP(Zn) bulk crystal reference sample shows, after surface annihilations at 0-7 keV, a flat curve which corresponds to annihilations of free positrons in the bulk. This level has been normalized to 1.0. Both the as-grown and annealed layers show between 0 and 7 keV a transition from the surface annihilations to annihilations in the layer. The flat part from 7 to 15 keV is characteristic of layer annihilations. Above 15 keV an increasing fraction of positrons start to penetrate into the substrate, causing the S parameter to decrease towards the bulk value.
The transition from surface annihilmions to layer annihilations can be used to estimate the effective positron diffusion length Leff. For both LT layers we get Leff ~-"200 A, which is much lower than the diffusion length of free positrons, L-~2000 A. The decrease in the diffusion length is due to trapping by defects and the trapping rate tc can be estimated [28] to be r = 100-300 ns- 1 in both layers. This is about two orders of magnitude higher than the free-positron annihilation rate, meaning that all positrons annihilate at defects. The value of S at 10-15 keV is thus characteristic of the defects (Sdef) trapping positrons in the layer. For the as-grown layer Sdef/Sbuik = 1.023, which is typical for monovacancies [30]. For the annealed layer the ratio SdJSbulk = 1.037 is tOO high for monovacancies and should correspond to divacancies or small vacancy clusters. The R parameter values of 1.0( 1 ) and 1.4(1) respectively in the two layers also support this conclusion. From preliminary electron irradiation experiments [30] we know that In vacancies produce parameters Svac/Sbulk and Wvac/Wbulk which are about the same as found in the present as-grown layer, whereas P vacancies induce much smaller changes in S and W. Thus we associate the vacancies with the In sublattice. According to calculations [31], the In vacancy is negative (2 - or 3 - ) when the Fermi level is above the midgap position. For a negative vacancy the trapping coefficient/~ is about 1015-1016 s-i. Using this value, we get the vacancy concentration [Vin]~ 1018-10 t9 cm- 3
P. Hautojiirvi et al. i
!
1.04
~ ~
Point defects in L T-MB E III- V materials
i
, _.~t
el
r, = s6o=c
4'• o /
1.03
/
1.0~ "t'l ,s-G,ow
~ 1.01 ,i...,q,. n
" I
1 oo
n
I
200
TEMPERATURE
n
Earlier [30] and present positron results suggest that it may originate from a vacancy or an ion-type defect. The effects of Be doping and variations in PH 3 flow during growth were also studied. The Be doping had no effect on the vacancies. The curves of as-grown layers with and without Be are superimposable. This is well understood if Be atoms are associated with oxygen [19]. An increase in the PH 3 flow from 8 to 22 sccm increased the vacancy signal slightly in the as-grown layers. It seems natural to have more In vacancies at higher P excess.
I
300 (K)
Fig. 4. Temperature dependence of S parameter measured at 12 keV positron energy in LT InP layer in as-grown state and after annealingat 560 °C. Figure 4 shows the S parameter at 10 keV as a function of the measurement temperature. One can see a very strong decrease with decreasing T. This is in contrast with the expected behaviour of /~ for a negative vacancy and means that at low temperature there are other ion-type negative centres which trap positrons at Rydberg states. The Rydberg states are shallow traps for positrons and the annihilation parameters are the same as in the bulk. When an increasing fraction of positrons annihilate at Rydberg states around negative ions, the S value decreases towards the bulk level. A rough analysis of the curves indicates that Ci.... =(0.1-0.5)Cvac, i.e. 1017-5 x 10 Is cm
21
- 3
It would be natural to identify these ions as Be atoms which are the intentional acceptor impurities. However, Be atoms have no effect on the carrier concentration--n = 5 x 1016 cm -3 with or without Be doping--most likely owing to their association with oxygen impurities [19]. This means that they are unable to trap positrons. Only residual impurities or intrinsic defects can be considered. Residual impurities at a level of 1017-1018 cm -3 are out of the question [19]. On the other hand, irradiation produces two types of positron traps, i.e. vacancies and negative ions [30]. This indicates that some intrinsic interstitial or antisite defects are acceptor type. It is likely that, as in the case of GaAs [26], the acceptors should have a double negative charge to trap positrons at Rydberg states. The most likely candidates are Inp antisites which, according to calculations [31], can be doubly negative when the Fermi level is above E c - 0.5 eV. It has been found previously from Hall measurements in n-type bulk InP that after heavy electron irradiation the Fermi level is pinned at an acceptor level located at Ec - 0.3 eV [32]. It is interesting that such a level is also observed in LT InP [20].
7. S u m m a r y
The As antisite is the well-identified native defect in LT GaAs observed earlier by EPR. As a midgap donor it plays the dominant role in the compensation. Positron annihilation gives information on acceptor-type defects, about which very little is known. Positrons detect vacancies in concentrations of 1018-1019 cm -3 in Be-doped layers grown at 225 °C. The vacancy concentration drops below 1016 cm -3 when the growth temperature is 300 °C. The vacancies are identified as V~a and so they can be the native acceptors compensating the native ASGa donors. On the other hand, multiple vacancies, divacancies or small clusters are observed in an undoped layer grown at 200 °C. No signal from Ga antisites is observed. The Pin antisite has recently been identified in LT InP by ODMR. It is believed to be responsible via auto-ionization for the high free-electron concentration always found in as-grown layers. Positron annihilation results show that as-grown InP layers contain high concentrations (1018-1019 cm -3) of vacancy defects identified as negative In vacancies. In addition, iontype acceptors, probably In antisites, in concentrations of 1017-1018 cm -3 are observed. The annealed layer contains small vacancy clusters.
Acknowledgments
We are grateful to J. H. von Bardeleben, J. Garcia, K. Evans and C. Stutz for providing us with samples.
References
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Point defects in L T MBE IH- V materials
4 P. Dreszer, W. M. Chen, K. Seendripu, J. A. Wolk, W. Walukievicz, B. W. Liang, C. W. Tu and E. R. Weber, Phys. Rev. B, 47(1993) 4111. 5 E Hautoj~irvi (ed.), Topics in Current Physics, Vol. 12, Positrons in Solids, Springer, Heidelberg, 1979. 6 L. Dorikens-Vanpraet, M. Dorikens and D. Segers (eds.), Positron Annihilation, World Scientific, Singapore, 1989. 7 P.J. Schultz and K. G. Lynn, Rev. Mod. Phys., 60 (1988) 701. C. Corbel, Scanning Microsc., 12(1987) 545. 8 K. M. Yu and Z. Liliental-Weber, Appl. Phys. Lett., 59 ( 1991 ) 3267. 9 H. J. von Bardeleben, M. O. Manaresh, D. C. Look, K. R. Evans and C. E. Stutz, Phys. Rev. B, 45(1992) 3372. 10 H.J. von Bardeleben, Y. Q. Jia, J. P. Hirzt, J. C. Garcia, M. O. Manaresh, C. E. Stutz and K. R. Evans, in ref. 1, p. 69. 11 M. O. Manaresh, D. C. Look, K. R. Evans and C. E. Stutz, Phys. Rev. B, 41 (1990) 10272. 12 M. Kaminska, E. R. Weber, Z. Liliental-Weber, R. Leon and Z. U. Rek, J. Vac. Sci. Technol. B, 7(1989) 710. 13 P. W. Yu, D. C. Reynolds and C. E. Stutz, Appl. Phys. Lett., 61 (1992) 1432. 14 S.A. McQuaid, R. C. Newman, M. Missous and S. O'Hagen, Appl. Phys. Lett., 61 (1992) 3008. 15 D. C. Look, Z.-Q. Fang and J. R. Sizelove, Phys. Rev. B, 47 (1993) 1441. 16 D. C. Look, D. C. Waiters, M. O. Manaresh, J. R. Sizelove, C. E. Stutz and K. R. Evans, Phys. Rev. B, 42(1990) 3578. 17 M. R. Melloch, N. Otsuka, J. M. Woodall, A. C. Warren and J. L. Freeouf, Appl. Phys. Lett., 57(1990) 1531. 18 A. C. Warren, J. M. Woodall, J. L. Freeout, D. Grischkowsky, M. R. Melloeh and N. Otsuka, Appl. Phys. Lett., 57(1990) 1331.
19 J. Ch. Oarcia, J. P. Hirtz, P. Maurel, H. J. yon Bardeleben and J. C. Bourgoin, in ref. 1, p. 277. 20 B.W. Liang, Y. He and C. W. Tu, in ref. 1, p. 283. B. W. Liang et al., Appl. l'hys. Lett., 60 (1992) 2104. 21 A. Claverie, J. Crestou and J. C. Garcia, Appl. Phys. Lett., 62 (1993) 1638. 22 J. M~ikinen, C. Corbel, E Hautoj~irvi, A. Vehanen and D. Mathiot, Phys. Rev., 42 (1990) 1750. 23 E. Soininen, J. M/ikinen, D. Beyer and E Hautoj~irvi, Phys. Rev. B, 46(1992) 13104, 24 M.J. Puska, C. Corbel and R. M. Nieminen, Phys. Rev. B, 41 (1990) 9980. 25 K. Saarinen, P. Hautoj~irvi, E Lanki and C. Corbel, Phys. Rev. B, 44(1991) 10585. C. Corbel, M. Stucky, P. Hautoj~irvi, K. Saarinen and E Moser, Phys. Rev. B, 38(1988)8192. 26 C. Corbel, E Pierre, K. Saarinen, E Hautoj~irvi and E Moser, Phys. Rev. B, 45(1991)3386. 27 S. Valkealahti and R. M. Nieminen, Appl. Phys. A, 35 (1984) 51. 28 K. Saarinen, P. Hautoj~irvi, J. Keinonen, E. Rauhala, J. Rfiis/inen and C. Corbel, Phys. Rev. B, 43(1991 ) 4249. 29 E. Soininen, J. M/ikinen, P. Hautoj~irvi, C. Corbel, A. Freundlieh and J. C. Grenet, Phys. Rev. B, 46(1992) 12394. 30 M. T6mqvist, J. Nissil~i, E Kiessling, C. Corbel, K. Saarinen, A. P. Seitsonen and E Hautoj~irvi, unpublished results, 1993. 31 M. Alatalo, R. M. Nieminen, M. J. Puska, A. P. Seitsonen and R. Virkkunen, Phys. R ev. B, 47(1993) 6381. A. E Seitsonen, R. Virkkunen, M. J. Puska and R. M. Nieminen, unpublished results, 1993. 32 V. N. Brudnyi, V. D. Kuznetsov, S. A. Vorobiev and A. A. Tsoi, Appl. Phys. A, 29 (1982) 219.
Materials Ai'ieHce am/t:)tgitwering, B22 1993 16 22
Point defects in III-V materials grown by molecular beam epitaxy at low temperature P. Hautojfirvi, J. Miikinen, S. P a l k o a n d K. S a a r i n e n
Laboratory of Physics, Helsinki Universityof Technology, SF-02150Espoo (Finland) C. C o r b e l a n d L. L i s z k a y
Centre d'Etudes NuclOairesde Saclay, INSTN, F-91191 Gif-sur-Yvette (France)
Abstract The present understanding of the point defects in GaAs and InP grown by molecular beam epitaxy at low temperature (LT) is briefly reviewed. New results on vacancies and ion-type acceptors obtained by positron annihilation are given. Depending on the growth temperature, Ga vacancies or small vacancy clusters are seen in LT GaAs in the concentration range l0 ts cm -3. No signal from Ga antisites is found. The LT InP layers contain vacancies, identified as In vacancies, in the concentration range 10 ~ cm 3. Ion-type acceptors, probably In antisites, are also seen in concentrations of l0 w cm-3. The annealed layer contains small vacancy clusters.
1. Introduction Low temperature (LT) growth of thin III-V layers has recently attracted much attention as a potential technique to obtain semi-insulating layers [1, 2]. Such layers are strongly needed for various micro- and optoelectronic device and integrated circuit applications. The layers contain point defects in extremely high concentrations. To identify and understand their role in the electrical and optical properties of the LT layers is a new challenge for various defect spectroscopic techniques. In LT GaAs and InP the main defects seem to be the group V antisite atoms AsGa and P~n, which have been identified as native donors [3, 4]. However, there is little information on native acceptors. Vacancies and antisites of group III atoms are generally believed to be acceptors. Positron annihilation is a unique method to identify vacancies [5-7]. In addition, as positive particles, positrons are also sensitive to ion-type acceptors. Here we present new information on the native acceptor defects in LT GaAs and InP by using positron annihilation. In this paper the main properties of the point defects in LT GaAs and InP are briefly summarized in Sections 2 and 3. The positron beam spectroscopy of thin layers is described in Section 4. The results on LT GaAs and InP in Sections 5 and 6 show that the layers contain high concentrations of native vacancies (V) tentatively identified as V(~a and VI,, respectively. LT (1921-5107/93/$6.00
InP layers also contain ion-type acceptors, most probably In antisites. The paper is summarized in Section 7.
2. Point defects in LT GaAs The defect properties of LT GaAs depend strongly on the growth temperature Tg, which varies from 200 to 300 °C. Analysis by particle-induced X-ray emission (PIXE) shows an As excess in the layer depending on the growth condition. The As excess is up to 1.5 at.% at Tg=190 °C and decreases monotonously as Tg increases. Layers grown at 300 °C or higher temperatures are stoichiometric within the sensitivity limit of 0.1 at.% of the PIXIE method [2, 8]. Studies by Rutherford backscattering and channelling indicate that As interstitial complexes rather than isolated As interstitials are present [8]. The only well-identified defect, however, is the arsenic antisite ASGa detected by the electron paramagnetic resonance (EPR) technique in its positive charge state ASGa÷ [3, 9]. Its concentration is 3-5 x 1018 cm -3 at Tg= 200 °C and decreases by an order of magnitude when Z. is increased to 300 °C [10]. The neutral state ASGafi is observed by optical absorption in the near-IR region [11]. Its concentration is a maximum of 2 × 10 2o c m -3 at Tg= 190 °C and decreases to 2 x 1019 cm -3 at Tg = 270 °C [2]. An open and controversial question is the photoquenching © 1993 - ELsevier Sequoia. All rights reserved
P. Hautojiirvi et al.
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Point defects in L T-MBE III-V materials
property of the AsGa defect in the LT layers, which seems to be different from that of the ASGa defect in semi-insulating bulk GaAs. Kaminska and Weber [2] have offered an interesting point of view. The strain in LT GaAs layers changes the local hyperfine field. The different properties of these two defects can be ascribed to differences in the local fields seen by the two defects rather than to differences in their nature. More generally, this question addresses the relation of the ASGa antisite to the EL2 centre in GaAs. Evidently ASGa is the main native donor in LT GaAs. It is activated at concentrations well above 10 ~8 cm -3. To compensate the ASGa+ centres, native acceptors are needed. Evidence of an intrinsic acceptor level at Ev + 0.3 eV has been given [12], which could be due to Ga vacancies. Yu et al. [13] have suggested that a complex seen by photolumiscence could be a VGa-As i pair. Measurements of the localized IR vibration modes of Si atoms suggest that Si-WGapairs are formed in Si-doped layers [14]. We show later in this paper that positron annihilation gives direct evidence of Ga vacancies at concentrations sufficiently high to account for the compensation. The conductivity of the as-grown layer is due to hopping within a defect band of ASGa antisites [15, 16]. After annealing of the layers at 600 °C, which always occurs in device applications, the normal conduction mechanism through thermally excited carders appears. The layers have become semi-insulating but the carrier lifetime and mobility are still low. The major annealing effect on the microstructure of the layers is the formation of As precipitates of size 6 nm to accommodate the As excess [17]. There is also a reduction in the concentration of ASGa° to below the detection limit of 1018 cm -3. In addition, the ASGa+ concentration drops to the level of 1017 cm-3 [9]. Two models have been suggested to explain the semi-insulating conductivity after annealing at 600 °C. (1) The point defect concentration is still high enough to explain the low carder lifetime and mobility. However, the question of the concentration of ASGa° and of other defects such as VGa remains open. (2) The high resistivity is due to metallic As precipitates which act as buried Schottky junctions. The space charge regions generated by the Schottky barriers overlap in the matrix, which is thus completely depleted of free carriers [18].
3. Point defects in LT InP
The success in the case of LT GaAs layers has stimulated efforts to grow LT InP and understand its electrical and structural properties. The growth temperature ranges from 130 °C upwards. The inter-
17
esting region is between 170 and 200 °C [19]. The amount of P excess is about 1 at.% at Tg= 170 °C and decreases with increasing growth temperature. The layers exhibit highly n-type conductivity except when they are passivated by hydrogen in the as-grown state [19, 20]. Beryllium doping does not compensate layers grown at Tg<300 °C [4, 19, 20]. Layers annealed at 600 °C stay conductive, although the P excess precipitates into small particles of size 3-7 nm [21 ]. The electrical properties of LT InP layers indicate that the dominant intrinsic defect is present in concentrations of 1019 cm -3 [20]. EPR measurements are not possible owing to the high conductivity of the layers. Recently Dreszer et al. [4] have found two deep donor levels at 0.11 eV above 0.23 eV below the conduction band. Optically detected magnetic resonance (ODMR) combined with photoluminescence reveals that the two levels are the first and second ionization levels of the P antisite double-donor defect. The upper level being resonant with the conduction band gives rise to the high free-electron via auto-ionization. To explain the photoluminescence spectrum, the authors need to call for an unidentified donor level at 0.8 eV above the valence band. We will see in Section 6 that in addition to the group V antisite atoms which seem to be the prevalent defects in LT InP, positron annihilation experiments indicate the presence of vacancies associated with the In sublattice in concentrations of 10 ~8-10 ~9 cm- 3
4. Positron annihilation technique 4.1. Positron dynamics in semiconductors
Energetic positrons in solids thermalize within a few picoseconds. In semiconductors the mean lifetime of free positrons is typically 250 ps. The motion and trapping of thermal positrons are in many respects analogous to those of free carriers apart from the "heavy" positron effective mass m*+ = 1.5m e. The positron diffusion coefficient at 300 K is around 2 cm 2 s- 1 [22, 23]. The diffusion is limited mainly by acoustic phonons, leading to the T -1/2 temperature dependence. Positrons feel electric fields and have a mobility of about 100 cm 2 V-1 s-1 t 300 K with a T -3/2 dependence. The Bloch state positron wavefunction is squeezed by the positron-ion repulsion into the interstitial regions between atoms. Thus an open-volume defect such as a vacant lattice cell is an attractive centre where positrons get trapped. Positron trapping is analogous to carder capture [24]. The trapping rate r is proportional to the defect concentration, r =/~ Cd. The charge state of a vacancy has a strong effect on the trapping coefficient/~. For a neutral vacancy/~ ~- 1014-1015 s-l
1N
P. Etautojdrvi et al.
/
Point defects in L T-MBE 111- V materials
independently of temperature. For a negative vacancy ~t ~ 1015-1016 S I and it increases by an order of magnitude at low temperatures. Positive vacancies do not trap positive positrons. Minimum vacancy concentrations required to trap positrons during their lifetime are about 10 l 5_ 1016 cm- 3. Owing to their different electronic environment, the annihilation characteristics of trapped positrons are different from those of free positrons in the bulk. The reduced electron density at a vacancy increases the positron lifetime by 10%-20%. Because of the smaller overlap of the trapped positrons with core electrons, the momentum distribution of the annihilating pairs becomes narrower. The change in the charge state of a defect causes measurable changes in the annihilation characteristics, thus giving the possibility to determine the ionization levels of the defect in the band gap [25]. Owing to coulombic interaction, a positron may also get trapped at weakly bound Rydberg states around negative ions. Thus information can also be obtained on ion-type acceptors [26].
4.2. Positron beam analysis Positrons from a radioactive source first hit a separate moderator crystal which has a negative work function for positrons. Thermalized positrons close to the moderator surface are spontaneously emitted into a vacuum. They form a beam which is accelerated and guided to the sample. In an experiment monoenergetic positrons are implanted at various depths in the sample by varying the beam energy E typically from () to 30 keV. In GaAs this energy range scans depths from 0 to 1.5 ~m. The positron-stopping profile is, however, rather wide [27], limiting the depth resolution of the technique. For reviews on positron beams see refs. 5 to 7. Since the start signal from the positron emission is not available in a beam, the positron lifetime cannot be measured, but the 511 keV line of the annihilation radiation is recorded. This line shows considerable Doppler broadening due to the motion of annihilating pairs. It is customary to use a line shape parameter S (or W ) defined as the ratio of the counts in the central (or wing) region of the annihilation line to the total number of counts in the peak. Owing to their high momentum, mainly core electrons contribute to the wing parameter W, whereas annihilations with the valence electrons fall predominantly in the energy region of the central parameter S. Both line shape parameters have characteristic values for each material, depending on the electron momentum distribution. When positrons are trapped at vacancytype defects, the annihilations with core electrons are reduced with respect to the perfect lattice, increasing the S parameter and decreasing the W parameter.
Hence the line shape parameters can be used to characterize vacancy defects in solids. The measured line shape parameter as a function of beam energy, S(E) (or W(E)), is a superposition of the characteristic values S~ for various annihilation states i. As an example, let us assume a defect layer in the nearsurface region of a sample. There are annihilations at the surface (surf), in the defects (def) and in the bulk (bulk) and we have
S(E) = F~urf(E) Ssurf~-/~\tcf(E) Sdef + Fbulk(E) Sbulk
(l)
Here b](E) denotes the fraction of annihilations in state i. To calculate the fractions Fi(E), the positron diffusion equation must be solved. Analysis of the experimental results is carried out by varying the positron-trapping profile ~c(x)=/zC~(x) in the diffusion equation until a satisfactory fit between the experimental and calculated S(E) is reached. The line shape parameters S and W depend on the fractions of positrons in various states. To overcome this difficulty, a new defect-specific parameter R is defined as [28, 29] R=
IS-&u,kl I W- Wbu,kI
(2)
If there are only two annihilation states, e.g. the bulk and a defect, R is independent of the fractions F~(E) and its value can be used to identify the defect.
5. Results on LT M B E GaAs
5.1. Samples The GaAs layers were grown in a Varian 360 system under normal, As-stabilized conditions at a growth rate of 0.8 /zm h-~ on semi-insulating GaAs substrates. A layer 4 /zm thick was grown at 200 °C. Two other layers of thickness 2.5 p m were grown at 225 and 300 °C and doped with Be to around 8 x 1019 cm -3. One sample was furnace annealed for 30 min in vacuum under face-to-face protection at successive temperatures of 300, 400 and 500 °C.
5.2. Results Figure 1 shows the experimental Doppler-broadening parameter S as a function of the positron beam energy E in three layers grown at 200, 225 and 300 °C. As reference sample we use a piece of semi-insulating LEC-grown bulk GaAs wafer which is known to contain no positron traps. The change in S(E) for E < 5 keV represents the transition from surface annihilations to annihilations in the layer. The flat part above
P. Hauto#irvi et al.
/
Point defects in L T-MBE III- V materials
MEAN STOPPING DEPTH (pro) 0.1 0.5 1.0 1.5 I
I
I
I
I
m
¢n 1.04
I
1.04
-o-¢
,,--
-
,--o--*•
1.03
=,,, = l =
225"C
o o-_.
~1.02
.
..
o_,_o
io.1 . 0 ~ LU
,t
#_o~~-
•
/
"P.
I
~ 1.01:Or)
I
To = 200"C
cc uJ
TG = 200"C
19
j300*C
I'~
,.ooll
--1.00
300"C
,
,
,
i
o 0
10 20 INCIDENT ENERGY (keV)
i
I 200
GaAs ref. i I 300
(K)
Fig. 2. Temperature dependence of S parameter measured at 15 keV positron energy in LT GaAs grown at various temperatures Tg.
cm-3. Above 5 keV implantation energy the positron trapping is saturated and the S = 1.037 value is characteristic of the defects in this layer. The defects are divacancies or small vacancy clusters [28, 29]. We performed furnace annealing of this layer up to 500 °C. No changes in the S ( E ) curves except at the surface (E < 2 keV) could be seen. This shows that the vacancy clusters are stable and difficult to anneal out. The high stability of vacancy clusters has also been observed earlier in ion-irradiated GaAs [28]. It is important to note that the total thickness of this layer is 4 #m, whereas positrons probe only the uppermost 1.5 # m region. Thus it is still open whether the region within the critical thickness (about 2 #m) from the substrate also contains vacancy clusters as the dominant defects. Figure 2 shows the temperature dependence of the S parameters in the layers measured at 10 keV positron energy. The message is clear: no temperature dependence. This means that the vacancy signal is not reduced at low temperature by positron trapping at Rydberg states around negative ions. This is in contrast to what has been observed in electron-irradiated GaAs, where in addition to Ga vacancies, negative-iontype acceptors identified as Ga antisites have also been observed [26]. We can conclude that the concentration of ion-type acceptors must be at least an order of magnitude less than that of vacancies. To determine whether the vacancies in the Tg-225 °C layer are in the Ga or As sublattice, we refer to their charge state. In semi-insulating GaAs, Ga vacancies are acceptors with a 2 - or 3 - charge, whereas As vacancies are positively charged [25]. Since positrons do not get trapped at positive defects, the traps can only be Ga vacancies. Such an identification a n d 1019
5 keV is characteristic of layer annihilations. To facilitate comparisons, the level of the S parameter in the reference sample has been normalized to 1.00. We can see a perfect overlap at the S parameter levels in the Tg=300 °C layer and the reference sample. This means that no vacancies in the layer are seen. Also, the transition from the surface value to the layer value extends up to 10 keV, indicating freepositron diffusion. The Tg= 225 °C has an S value of 1.023. This is well above the reference bulk level, which indicates the presence of vacancy-type defects in the layer. The very rapid increase from the surface S value to the plateau means that the positron diffusion length L + ~ 100/k is very short compared with the free-positron diffusion length L+ ~ 1600 A [23]. This indicates a high concentration (about 1019 cm-3) of defects trapping positrons, which means that all positrons between 3 and 20 keV annihilate when trapped. The S value of 1.023 is characteristic of the defects in the layer. This value is close to those found earlier for As and Ga monovacancies. Thus we can conclude that the dominant traps are monovacancies in this layer. The decrease in S above 20 keV is due to the increasing fraction of positrons, which at higher energies penetrate into the substrate with a negligible defect concentration compared with the overlayer. The S value in the Tg= 200 °C layer is 1.040. This is even higher than the value for monovacancies. The increase in S between 0 and 5 keV gives an effective positron diffusion length L+ = 1 7 0 /k. The corresponding positron trap concentration is between 10 ]8
I I00
TEMPERATURE
30
Fig. 1. Line shape parameter S as a function of positron beam energy in LT GaAs layers grown at various temperatures T~. The Tg--300 °C curve coincides with that of the semi-insulating GaAs crystal (not shown) used as reference.
o.~o__~:~o-g-
il,~-,:W
20
P. Hautojdrvi et al.
/
Point defects in L T-MBL 111- V materials
leads us to conclude that the vacancies act as native acceptors. It is interesting that in the layer grown at 300 °C no Ga vacancies are seen. This means that the concentration of Ga vacancies is less than 101~' cm -3. Von Bardeleben et al. have measured the concentration of ASGa+ as a function of growth temperature. Their results show the same trend as for our VG.: when Tg is 300°C or higher, the AsG~+ concentration has decreased by more than one order of magnitude from the layers grown at 200 °C [10]. In conclusion, positron annihilation shows the existence of vacancies and vacancy clusters in concentrations of 1018-1019 cm -3 in the Be-doped layer grown at 225 °C and in the undoped layer grown at 200 °C respectively. No signal from Ga antisites is observed. The vacancies are associated with V~a and so they can be the acceptors compensating the native Asoa donors observed earlier by EPR. In the Be-doped layer grown at 300 °C the V ~ concentration is below 1016 cm-3.
M E A N S T O P P I N G DEPTH (/am) 0.1 0.5 1.0 1.5 I
I
l
I
~oO°%°°°°°-~°Oo~o
m 1.04 nuJ
o
~ 1.03 <
°°- TA -- 5600C O
~ 1.02 w •
"I" 1 . 0 1
~
AS-GROWN
w Z
~ 1.00
InP ref. IA
0
I
I
10
20
=
I
30
INCIDENT E N E R G Y (keV)
Fig. 3. Line shape parameter S as a function of positron beam
energy in LT InP layers in as-grown state and after annealingat 560 °C. As reference sample an [nP (Zn:2 × 10~s cm -3) crystal has been used.
6. Results on LT MBE InP
6.1. Samples Several InP layers were grown by gas source molecular beam epitaxy at Thomson, Corbeville. For details of the growth system see ref. 19. The growth temperature was 180 °C and the thickness of the layers 1.5/~m. The PH 3 flow rate was 8 or 22 sccm (standard cm 3 min- 1). Most of the layers were doped with Be to 5 x 1017 cm-3. All the layers were n type with a carrier concentration of 5 x l0 t6 cm -3 even after annealing at 560 °C for 30 min. As a "defect-free" reference we used a Czochralski-grown InP (Zn: 2 x 10 Is cm-3) bulk wafer which, according to positron lifetime measurements, does not exhibit any positron trapping by defects.
6.2. Results Figure 3 shows the experimental Doppler-broadening parameter S as a function of the positron beam energy E for three different samples. The InP(Zn) bulk crystal reference sample shows, after surface annihilations at 0-7 keV, a flat curve which corresponds to annihilations of free positrons in the bulk. This level has been normalized to 1.0. Both the as-grown and annealed layers show between 0 and 7 keV a transition from the surface annihilations to annihilations in the layer. The flat part from 7 to 15 keV is characteristic of layer annihilations. Above 15 keV an increasing fraction of positrons start to penetrate into the substrate, causing the S parameter to decrease towards the bulk value.
The transition from surface annihilmions to layer annihilations can be used to estimate the effective positron diffusion length Leff. For both LT layers we get Leff ~-"200 A, which is much lower than the diffusion length of free positrons, L-~2000 A. The decrease in the diffusion length is due to trapping by defects and the trapping rate tc can be estimated [28] to be r = 100-300 ns- 1 in both layers. This is about two orders of magnitude higher than the free-positron annihilation rate, meaning that all positrons annihilate at defects. The value of S at 10-15 keV is thus characteristic of the defects (Sdef) trapping positrons in the layer. For the as-grown layer Sdef/Sbuik = 1.023, which is typical for monovacancies [30]. For the annealed layer the ratio SdJSbulk = 1.037 is tOO high for monovacancies and should correspond to divacancies or small vacancy clusters. The R parameter values of 1.0( 1 ) and 1.4(1) respectively in the two layers also support this conclusion. From preliminary electron irradiation experiments [30] we know that In vacancies produce parameters Svac/Sbulk and Wvac/Wbulk which are about the same as found in the present as-grown layer, whereas P vacancies induce much smaller changes in S and W. Thus we associate the vacancies with the In sublattice. According to calculations [31], the In vacancy is negative (2 - or 3 - ) when the Fermi level is above the midgap position. For a negative vacancy the trapping coefficient/~ is about 1015-1016 s-i. Using this value, we get the vacancy concentration [Vin]~ 1018-10 t9 cm- 3
P. Hautojiirvi et al. i
!
1.04
~ ~
Point defects in L T-MB E III- V materials
i
, _.~t
el
r, = s6o=c
4'• o /
1.03
/
1.0~ "t'l ,s-G,ow
~ 1.01 ,i...,q,. n
" I
1 oo
n
I
200
TEMPERATURE
n
Earlier [30] and present positron results suggest that it may originate from a vacancy or an ion-type defect. The effects of Be doping and variations in PH 3 flow during growth were also studied. The Be doping had no effect on the vacancies. The curves of as-grown layers with and without Be are superimposable. This is well understood if Be atoms are associated with oxygen [19]. An increase in the PH 3 flow from 8 to 22 sccm increased the vacancy signal slightly in the as-grown layers. It seems natural to have more In vacancies at higher P excess.
I
300 (K)
Fig. 4. Temperature dependence of S parameter measured at 12 keV positron energy in LT InP layer in as-grown state and after annealingat 560 °C. Figure 4 shows the S parameter at 10 keV as a function of the measurement temperature. One can see a very strong decrease with decreasing T. This is in contrast with the expected behaviour of /~ for a negative vacancy and means that at low temperature there are other ion-type negative centres which trap positrons at Rydberg states. The Rydberg states are shallow traps for positrons and the annihilation parameters are the same as in the bulk. When an increasing fraction of positrons annihilate at Rydberg states around negative ions, the S value decreases towards the bulk level. A rough analysis of the curves indicates that Ci.... =(0.1-0.5)Cvac, i.e. 1017-5 x 10 Is cm
21
- 3
It would be natural to identify these ions as Be atoms which are the intentional acceptor impurities. However, Be atoms have no effect on the carrier concentration--n = 5 x 1016 cm -3 with or without Be doping--most likely owing to their association with oxygen impurities [19]. This means that they are unable to trap positrons. Only residual impurities or intrinsic defects can be considered. Residual impurities at a level of 1017-1018 cm -3 are out of the question [19]. On the other hand, irradiation produces two types of positron traps, i.e. vacancies and negative ions [30]. This indicates that some intrinsic interstitial or antisite defects are acceptor type. It is likely that, as in the case of GaAs [26], the acceptors should have a double negative charge to trap positrons at Rydberg states. The most likely candidates are Inp antisites which, according to calculations [31], can be doubly negative when the Fermi level is above E c - 0.5 eV. It has been found previously from Hall measurements in n-type bulk InP that after heavy electron irradiation the Fermi level is pinned at an acceptor level located at Ec - 0.3 eV [32]. It is interesting that such a level is also observed in LT InP [20].
7. S u m m a r y
The As antisite is the well-identified native defect in LT GaAs observed earlier by EPR. As a midgap donor it plays the dominant role in the compensation. Positron annihilation gives information on acceptor-type defects, about which very little is known. Positrons detect vacancies in concentrations of 1018-1019 cm -3 in Be-doped layers grown at 225 °C. The vacancy concentration drops below 1016 cm -3 when the growth temperature is 300 °C. The vacancies are identified as V~a and so they can be the native acceptors compensating the native ASGa donors. On the other hand, multiple vacancies, divacancies or small clusters are observed in an undoped layer grown at 200 °C. No signal from Ga antisites is observed. The Pin antisite has recently been identified in LT InP by ODMR. It is believed to be responsible via auto-ionization for the high free-electron concentration always found in as-grown layers. Positron annihilation results show that as-grown InP layers contain high concentrations (1018-1019 cm -3) of vacancy defects identified as negative In vacancies. In addition, iontype acceptors, probably In antisites, in concentrations of 1017-1018 cm -3 are observed. The annealed layer contains small vacancy clusters.
Acknowledgments
We are grateful to J. H. von Bardeleben, J. Garcia, K. Evans and C. Stutz for providing us with samples.
References
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