- Email: [email protected]
GaAs MQWs for advanced photovoltaic devices
ARTICLE IN PRESS Journal of Crystal Growth 311 (2009) 4293–4300
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Structural and optical characterization of MOVPE grown InGaP/GaAs MQWs for advanced photovoltaic devices M. Longo a,1, A. Parisini b,, L. Tarricone b, S. Vantaggio b, C. Bocchi c, F. Germini c, L. Lazzarini c a
CNR-INFM, Dipartimento di Fisica, Universita di Parma, Viale G.P. Usberti, 7/A, 43100 Parma, Italy CNISM, Dipartimento di Fisica, Universita di Parma, Viale G.P. Usberti, 7/A, 43100 Parma, Italy c CNR-IMEM Parco Area delle Scienze 37/A, 43010 Loc. Fontanini-Parma, Italy b
a r t i c l e in fo
abstract
Article history: Received 30 January 2009 Received in revised form 12 June 2009 Accepted 5 July 2009 Communicated by K.W. Benz Available online 16 July 2009
The optimization of the electronic properties of InGaP/GaAs MQWs, to be inserted in multilayers heterostructure for novel photovoltaic devices, was performed by structural, optical and photoelectrical measurements. Different sequences of nominally undoped InGaP and GaAs alternated layers were grown by low-pressure metalorganic vapour phase epitaxy, employing tertiarybutylarsine and tertiarybutylphosphine as metalorganic precursors for the V-group elements. In order to minimize the As/P exchange effect, the interface In segregation, and to control the whole lattice matching, single and multi-quantum wells (MQWs) with different: (i) periods, (ii) well widths, (iii) growth temperatures, (iv) gas-switching sequences at the interfaces and (v) indium concentrations in the InGaP alloy, were prepared and investigated. The interface sharpness and the compositional fluctuation of thick MQW region containing up to 40 well-barrier sequences were investigated for the modelling, realization and evaluation of test structures based on low-dimensional systems for third generation solar cells. & 2009 Elsevier B.V. All rights reserved.
PACS: 81.05.Ea 81.15.Gh 81.07.St Keywords: A1. InGaP/GaAs interfaces A3. Gas-switching sequence A3. MOVPE A3. Multi-quantum wells B3. Solar cells B2. Semiconducting indium gallium phosphide
1. Introduction The lattice-matched InGaP/GaAs heterostructure is generally recognized as an interesting system for microelectronic and optoelectronic applications. In particular, it can be inserted in a multilayer structure to obtain a better matching of the solar spectrum, and then enhance the performances of photovoltaic devices. The quantum well solar cells (QWSC) technology, in fact, appears to be a potential candidate to achieve high conversion efficiency devices [1]. The crucial technological requirements for the modelling, realization and test of prototypal photovoltaic devices based on low-dimensional systems, include: (i) as sharp as possible superlattice interfaces, (ii) control of residual lattice strain and structural disorder, (iii) high regularity of the QW periodicity and good correspondence between modelling and Corresponding author. Tel.: +39 0521 905272; fax: +39 0521 905223.
E-mail address: [email protected] (A. Parisini). Present address: Laboratorio Nazionale MDM, Via C. Olivetti, 2, 20041 Agrate Brianza (MI), Italy. 1
0022-0248/$ - see front matter & 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2009.07.015
experimental determination of the bandgap profile, (iv) reduction of background impurity and deep level density. Thus, in the present work, an assessment of the MOVPE growth conditions was carried out to realize QW-based test structures satisfying the (i)–(iv) requirements, to evaluate their potential performances as quantum photovoltaic devices. The advantages of using InGaP with respect to AlGaAs, beyond the possibility to realize Al-free devices, are mainly related to: (i) lower oxidation tendency, (ii) higher mechanical stability and (iii) lower interfacial recombination velocity. On the other hand, when grown by metalorganic vapour phase epitaxy (MOVPE), the InGaP/ GaAs system suffers some problems at the interfaces, such as In memory effects or As/P intermixing. In particular, Indium memory effects and surface segregation can lead to the formation of unintentional InxGa1xAsyP1y layers at the interfaces. Such uncontrolled quaternary interlayer, evidenced by a low-energy undesired optical emission, can affect the lattice match composition and the interfacial morphology, resulting in a degradation of the device properties. Various attempts to improve the interface sharpness of the InGaP/GaAs heterojunction, by exploiting
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different gas switch sequences (GSSs) or by interposing proper interlayers (ILs), have been previously carried out [2–7], the argument being again object of debated research. With regard to safety and purity growth issues, one interesting option is the use of alternative precursors tertiarybutylarsine (TBAs) and tertiarybutylphosphine (TBP), in which one of arsine (phosphine) hydrogen functions is replaced by a heavier butylgroup. In such a way both the toxicity, the oxygen incorporation and the background doping should be significantly lowered by virtue of TBAs and TBP low dissociation temperatures, which also allow the growth below Tg ¼ 600 1C [8,9]. Indeed, (InGaAsP)-based devices, grown by using TBAs and TBP and showing at least comparable otherwise better performances, have been recently realized [10–13]. By employing a MOVPE reactor using TBAs and TBP precursors, we recently investigated the effects of different gas-switching sequences for the growth of multi-quantum wells (MQWs) based on the lattice-matched system InGaP/GaAs [14]. It resulted that the insertion of a few-monolayer thick GaAsP IL at the direct GaAs-on-InGaP interface and the use of proper growth interruption purges lead to optimum optical properties of the structures, by suppressing the undesired anomalous low-energy emission due to In memory effect. Moreover, the emission energies due to transitions between confined quantum states, measured as a function of the well thickness, were in good agreement with the theoretical expectation. In this work, the effects due to the IL insertion and the assessment of growth conditions suitable for the control of the residual strain in the structure are accurately discussed, by taking into account previous results reported in Refs. [14,15] to give a complete view of the work. Optical and photoelectrical investigation on p+-MQW-n test structure are reported, to demonstrate the feasibility of MQW structures with a consistent number of periods, as required for advanced photovoltaic applications [16].
2. Experimentals The growth of the InGaP/GaAs heterostructures was performed by a horizontal low-pressure MOVPE Aixtron reactor (AIX200 RD). In order to optimize the interfaces, previously found conditions were applied [14]. Trimethylgallium (TMGa), trimethylindium, TBAs and TBP metalorganic precursors were used for III–V-group
elements, while hydrogen-diluted disilane was used for n-doping and dimethylzinc for p-doping. Palladium purified H2 was used as a carrier gas. The substrates were 2 inch GaAs (0 0 1) exact wafers, pre-treated by standard etching procedures. The TMGa partial pressure was fixed at 5.6 103 mbar and the V/III ratio was 25.5 for GaAs and 14 for InGaP, the growth was performed at a temperature of 560–600 1C, a total pressure of 50 mbar and a total gas flow of 6.8 l/min. The typical growth rate turned out to be of 1.2 mm/h for GaAs and 2.2 mm/h for InGaP. A particular gasswitching sequence was employed to optimize the optical and morphological quality of the GaAs–InGaP interfaces, as in Ref. [14]. Different sequences of unintentionally doped and nominally pseudomorphic InGaP/GaAs MQWs were grown at Tg ¼ 600 1C, confined within InGaP cladding layers. Single QWs, as well as MQWs, with different periods, variable well size, gas-switching sequence at the interfaces and indium concentration in the InGaP alloy were prepared. In all cases a 200–300-nm-thick GaAs buffer layer was grown on top of the substrate. GaAs wells ranging in thickness between 3 and 10 nm and 80–150 nm thick InGaP barriers were then deposited. The main growth parameters of the samples are collected in Table 1. A background p-type doping of the InGaP layers resulted by electrical investigation of the samples [17,18]. To allow an easier correlation between the results of the present work and those reported in our previous publications for samples featured by same structure and deposition conditions, we point out that the new simple labels here used substitute the identification labels of past publications as follows: samples 1, 2, y stand for samples IGPMQW1, IGPMQW2, y; C1 stands for sample IGPCELL1. The grown structures were investigated through structural and optical measurements. The structural analysis was performed by X-ray diffraction (XRD) and electron microscopy methods. A Philips high-resolution X-ray diffractometer, equipped with a graded multilayer mirror for obtaining a quasi-monochromatic and -parallel beam of high intensity, and a Bartels channel-cut monochromator for selecting the CuKa1 radiation, was used for XRD measurements. To reduce the diffuse scattering coming from the sample, a thin slit was placed in front of the detector with an acceptance angle nearly 0.51 and the diffraction profiles were recorded in the o2y scan mode, by using the 004 Bragg reflection. This arrangement extended the range of the measured
Table 1 Main features of the grown InGaP/GaAs heterostructures. MQW
GaAs well width Nominal (nm)
Measured by TEM (nm)
13
8 8 2 4 6 4 4 4 8
8.370.5 – 3.570.5 5.570.5 7.270.5 5.770.5 4.970.5 4.670.5 –
14
8
–
17 18
8 2 (top) 4 6 8 (bottom) 8
8.670.5 3.071.0 5.070.5 6.570.5 8.570.5 –
1 2 5
6
C1
(top) (middle) (bottom) (top) (middle) (bottom)
GaAsP interlayer (IL)
QW periods
XRD measured indium molar fraction in InGaP
Absent (no IL) IL at direct interface (4 s) Same IL at direct interfaces (4 s)
20 1 3
0.48 – –
Different ILs at direct interfaces (5-3-1 s)
3
0.47
IL at direct+reduced GIT at indirect interface (2 s) IL at direct (4s)+IL at indirect (4 s) interfaces IL at direct interfaces IL at direct interfaces
1
–
1
–
40 20
0.46 0.445
IL at direct interfaces
30
0.488–0.495 (analysis on the whole grown area)
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3. General assessment of MQW structures The milestone of this work was the assessment of the growth parameters for the realization of structures formed by a sequence of several QWs, according to the fundamental requirements for a low-dimensional system-based device. In this section, we recall–giving new details–the main results previously obtained [14,15] and complete the study by adding new experimental data on the final test structure. The achievement of a good morphological quality and a very low structural and compositional disorder had been already confirmed by the DF-TEM image and the XRD curve reported in Ref. [15], resulting in an excellent reproducibility over tens of MQWs. The investigation concerned a MQW heterostructure (sample 1) grown without any particular gas-switching sequence at the interfaces, therefore, without any intentional GaAsP interlayer. The nominal well and barrier thickness of 8 and 12 nm, respectively, show a satisfying agreement with respect to the values of 8.3 and 9.5 nm, measured by TEM images. In this structure, where 20 InGaP/GaAs periods are stacked, the interfaces exhibit a roughness around a few nm, with a lateral period of about 100 nm, slightly increasing from the bottom to the top of the MQWs. Such a lateral unhomogeneity of the interface is also evidenced by a shoulder in the FFT–PL spectrum reported in Ref. [14] (where, the sample is labelled IGP1). However, a shift towards anomalously low values of the PL emission of the MQW reveals the formation of a quaternary InGaAsP layer at the direct GaAs-on-InGaP interface [14].
1 (004)
0
calculated experimental
-1
-2 log (reflectivity)
intensity up to 106 times the value of the substrate Bragg peak. The experimental XRD curves were analysed by means of a best fit procedure based on the X-ray diffraction dynamical theory for distorted crystals [19,20]. The layer thickness t, the static Debye–Waller factor (DW) accounting for the lattice disorder and the lattice spacing modification Dd/d, were the parameters of the calculations. The best fit was obtained by means of a least-squares method for the difference (w) between the experimental and calculated curves. Transmission electron microscopy (TEM) investigations were carried out on /11 0S cross-section specimens by using the (2 0 0) bright field (BF) and dark field (DF) imaging mode, the latter being sensitive to the chemical composition. In this imaging condition, the width of the GaAs quantum wells was measured, by means of proper image analysis software, and compared with the values obtained by HRXRD. Low-temperature cathodoluminescence (CL) investigations were performed on a Cambridge Stereoscan 360 SEM equipped with an Oxford Instruments mono-CL system, with the specimen observed in plan view. Photoluminescence (PL) spectra were detected through a Bruker-IFS66-FFT spectrometer, using the 488 nm line of a 5 W Ar+ ion laser and the 632.8 nm line of a He–Ne laser as excitation and reference light, respectively. A spectral resolution of 0.1 meV (about 1 cm1) was obtained in the investigated spectral range. Finally, photocurrent spectra were obtained through measurements in the temperature range (30–300 K) on p–i–n like structures, whose intrinsic region was formed by a sequence of InGaP/GaAs QWs. The junctions were illuminated by a halogen lamp (100 W), focused into the entrance of a Spex monochromator (0.5 m focal length), with a spectral resolution better than 2 meV. The spectral responses of the MQW p–i–n diodes were obtained by preparing on the top surface a pattern of mesa structures, 500 mm in diameter, by conventional photolithography, as described in Ref. [18]. The PR signal was detected by conventional lock-in techniques and normalized to the system spectral response by using a calibrated Si photodetector.
4295
-3
-4
-5
-6
-7 -3000
-2000
-1000 0 1000 Δθ (arcsec)
2000
3000
Fig. 1. XRD 004 experimental (dot) and calculated (line) profiles of a single QW structure (sample 2) in which a GaAsP interlayer (IL) was inserted at the direct interface.
The efforts to improve the interface quality were subsequently focused on single QWs, where sharp interfaces were obtained when a GaAsP interlayer (IL) is inserted at the direct GaAs-onInGaP interface [14,15]. This expedient was expected to reduce the diffusion/segregation of In through the interface. Fig. 1 shows the XRD spectrum of sample 2, superimposed to the theoretical fit, by which a 1.7 nm thickness was obtained for the GaAsP interlayer, comparable to the nominal value of 1.5 nm. The fair agreement between experiment and theory gives evidence in favour of the high quality of the structure and the good control of the growth parameters. In addition, the PL emission energy of such structure approaches the expectation for the optical transition between the n ¼ 1 heavy hole subband and the n ¼ 1 electron subband in the wells (hh1e1 transition), once the 2D exciton binding energy is taken into account [14]. The effect of either a variation of the growth interruption time (GIT) in sample 13 or the insertion of a GaAsP IL in sample 14 at the indirect interface were also investigated, without varying the well and barrier thickness. As it is shown in Fig. 2, in both cases the sharpness of the interface degrades with respect to sample 2, in fact, low-energy peaks appears again in the PL spectrum of samples 13 and 14, also in this case attributable to the formation of an undesired quaternary alloy [21]. In addition, the main emission peak, detected, as expected, at about 1.535 eV in samples 13 and 2, shifts toward higher energies (1.557 eV) in sample 14. This peak is due to the hh1e1 excitonic recombination, whose energy position depends on the effective well width, therefore, an apparent reduction of the QW thickness from 8 to 5 nm seems to be induced in the QW of sample 14. 3.1. Optimization of the GaAsP IL thickness at the direct interface With the aim to more deeply investigate the role of the GaAsP IL, two samples containing a different sequence of three stacked
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average over the three wells, showing a single emission peak, with a minor effect on the linewidth. The active role of the IL in limiting the As/P intermixing and the In diffusion was, however, confirmed by XRD analysis on sample 6 (Fig. 4). Also, a good control of the lattice match at the InGaP/GaAs interfaces resulted again, with /xInS ¼ 0.47 for each barrier. No effects, ascribable to the interface broadening were observed, the roughness being in any case again limited within 1 nm, with the upper (indirect) interface
Fig. 2. FFT–PL spectrum recorded at 12 K from MQW structures grown by different gas-switching sequences (GSSs) at the indirect interfaces. Arrows indicate in each spectrum the hh1e1 excitonic transition.
Fig. 3. CL spectra of sample 6, recorded at different probing depths (electron accelerating voltages). The peaks are relative to the hh1e1 excitonic transition.
1 fitting experimental 0
(004)
-1
log (reflectivity)
InGaP/GaAs QWs were grown, that is samples 5 and 6, having either the same GaAsP IL thickness (expected thickness around 1.7 nm). Sample 5 contains three single QWs of different widths (6, 4 and 2 nm, from the bottom to the top, respectively), sample 6 contains three single QWs of nominally identical width (4 nm), but different GaAsP ILs thickness (0.4, 1.2 and 2.0 nm, from the bottom to the top, respectively). In both samples the barriers were 150 nm thick. The well-width values obtained by TEM in sample 5, where the IL was four monolayers thick, were found to be slightly higher than the nominal ones, that is 7.270.5, 5.570.5 and 3.570.5 nm, respectively. The low-temperature PL spectrum of this sample, already shown in Fig. 2b of Ref. [14], exhibits three well-separated peaks, lying at 1.545, 1.571 and 1.620 eV, respectively, which quite well follow the theoretical trend expected for the hh1e1 radiative transition. The slight but systematic red shift of the optical emission data with respect to the theoretical energies, in spite of the inclusion of the 2D exciton binding energy, will be briefly discussed later. The TEM data seem to suggest that the presence of the IL only slightly influences the well width. However, the investigation of sample 6 revealed that, when the IL thickness increases to 2 nm, the morphology of the well degrades and the effective well width increases from a nominal value of 4 nm to an effective value of 5.771 nm, significantly more with respect to the other two quantum wells, which appear 4.970.5 and 4.670.5 nm wide, for a 1.2 and 0.4-nm-thick IL, respectively [14]. This result was also confirmed by space-resolved cathodoluminescence (CL) measurements. By taking advantage of the dependence of the beam penetration depth on the kinetic energy of the exciting electrons, and with the aid of Montecarlo simulations (not shown here), we were able to selectively investigate the different wells through the variation of the voltage accelerating the incident electrons. As it appears from Fig. 3, the 3 kV CL spectrum, obtained for sample 6 from the top well prevalently, shows a main emission peak (E ¼ 1.567 eV) at a slightly lower energy with respect to the energies of the emission peaks dominating the CL spectrum, both at an intermediate accelerating voltage (V ¼ 4 kV, emission from the middle QW) and at the highest accelerating voltage (V ¼ 5 kV, emission from the bottom QW). This behaviour suggests a larger width of the top well with respect to the other two. Nevertheless, such a spread is too limited to be evidenced by PL measurement, which gives an
-2
-3
-4
-5
-6 -2000
-1000
0 Δθ [arcsec]
1000
2000
Fig. 4. Experimental (dot) and calculated (line) 004 X-ray diffraction profiles of sample 6.
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(ii)
2.0 (iii)
PL intensity (a.u.)
T = 10 K 1.5
(iv)
1.0
0.5
0.0 1.2
1.3
1.4
1.5 Energy (eV)
1.6
1.7
(i) Lw = 2 nm
(ii) Lw = 4 nm
(iii) Lw = 6 nm
(iv) Lw = 8 nm
Fig. 5. PL peak related to the hh1e1 excitonic transition at T ¼ 10 K (a) and BF TEM image (b) acquired from a lattice mismatched structure containing 40 GaAs quantum wells, 8 nm thick (sample 17).
Fig. 6. PL spectrum at T ¼ 10 K (a) and (2 0 0) DF TEM image (b) of a lattice mismatched structure containing four series of quantum wells, each containing four periods, having well width decreasing from 8 to 2 nm (sample 18). Only the hh1e1 excitonic PL peaks relative to the MQWs having well width, respectively, of (ii) 4 nm, (iii) 6 nm and (iv) 8 nm are detected. In the image (b), the corresponding sequence of MQWs is evidenced.
of the wells worse than the lower (direct) one, and no lattice defect due to dislocations being evidenced [14,15].
3.2. Lattice mismatch effects The requirement of growing heterostructures containing a consistent number of QW periods imposes an accurate control of the InGaP/GaAs lattice mismatch and the effects of the residual strain. Different MQW structures were grown, having different stacking sequences and well widths. In particular, we now discuss the results relative to a series of 40 QWs, having the same well width Lw ¼ 8 nm (sample 17), and a structure, sample 18, formed by four groups of MQWs (of four QWs each), having Lw decreasing from 8 nm (in the bottom QWs), to 6, 4 and 2 nm (in the top QWs). At each direct interface, a 1.7-nm-thick GaAsP IL was interposed. Fig. 5a and b, respectively, show the low-temperature FFT–PL spectrum and the TEM image obtained for sample 17, whose In mole fraction resulted intentionally low, xIn ¼ 0.46, with respect to the perfect lattice matching value of xIn ¼ 0.484. The main peak in the PL spectrum at E ¼ 1.531 eV, with a FWHM of 0.013 eV, is consistent with the expectations, while the weak shoulder
observed at E ¼ 1.46 eV could be attributed to interface defects or non-intentional impurities. The TEM image shows a strong modification of the MQW periodicity after about 70 interfaces, due to the elastic strain relaxation. Because of the residual lattice mismatch, the structure is locally perturbed but has not relaxed as a whole and misfit dislocations have not been detected at any interface. A similar analysis on sample 18 reveals three PL emission peaks due to the first three sets of QWs (Fig. 6a), at E ¼ 1.533, 1.546 and 1.563 eV from the bottom to the top of QWs set, whereas the 2 nm wide QW is completely destroyed by the lattice relaxation (Fig. 6b), which in this structure appears after 30 interfaces, again because of a low In content (In mole fraction ¼ 0.445). The elastic energy accumulated in the structure partially favours the interface roughness, in fact, in the literature the latter has been considered as the main cause of the degradation of the top layers in InGaP/GaAs MQWs. Indeed, this effect has been evidenced also in lattice-matched structures containing many interfaces [22].
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Table 2 Scheme of the optimized structure C1, containing 30 staked MQWs embedded in doped InGaP cladding layers. Layer
Material
Thickness (nm)
Dopant type
Dopant
Dopant concentration (cm3)
Cap Top cladding layer Top spacer layer Barrier Well Direct interface IL Bottom spacer layer Bottom cladding layer Buffer Substrate
GaAs InGaP InGaP InGaP GaAs GaAsP InGaP InGaP GaAs GaAs
10 300 50 12 8 1.5 100 300 250
Zn Zn
p++ p+
2 1018 1018
Si Si Si
n+ n+ n+
2 1018 2 1018
30 periods
Fig. 7. FFT–PL spectrum at T ¼ 10 K obtained from the optimized p–i–n structure containing a MQW of 30 period (sample C1).
4. Study of p+-MQW-n+ test structures The assessment of the growth parameters was exploited for the design and realization of a p–i–n test structures containing MQWs in the intrinsic region (sample C1), which is the basic structure for the realization of novel photovoltaic devices (QWSCs). Thirty lattice-matched InGaP/GaAs QWs were stacked. In order to limit impurity diffusion from the cladding layers into the MQW region, in this case, spacer layers were properly interposed. The details of the structure are reported in Table 2. A very high optical quality of the sample appears from the low-temperature (12 K) FFT–PL spectrum, dominated by the narrow emission from the Lw ¼ 8 nm width wells (Fig. 7). The PL peak, due to the hh1e1 transition, lies at E ¼ 1.530 eV and has a full-width at half-maximum (FWHM) of 12 meV. The energy peak position was also confirmed by a CL investigation, indicating a good reproducibility over all the samples prepared with the same modelling parameters. The optical properties of the test structures were more extensively investigated through photocurrent spectroscopy (PCS), which gives information complementary to PL and CL. In fact, whereas the latter techniques reveal the more probable optical recombination process (emission of photons), in the PCS, photons of variable wavelength are absorbed, giving information on permitted excitation transitions. The PCS investigation of MQWs, in particular, evidences the effects of spatial confinement
of carriers and offers the possibility to compare nominal and effective QW sizes [23]. In addition, it is strongly influenced by the collection efficiency of carriers at the contacts, then by the transport mechanisms of photo-generated carriers through the QWs. Fig. 8a shows the photocurrent spectrum of sample C1, recorded at T ¼ 300 K. The spectrum, extends between the values of the GaAs and InGaP energy gaps (window effect), as expected. The shoulders observed at low energies are related to the hh1e1 and lh1e1 excitonic transitions; generally, the evidence of the splitting between heavy and light hole related peaks is a fingerprint of an optimum structural quality of the MQW, since a high density of defects and the internal field due to the presence of strain generally broaden the optical features, resulting in an enlarged lineshape. Some additional information on the structural quality of the samples can be obtained by the analysis of the whole PC spectrum at different temperatures. As a minor feature, we mention the evidence of a negligible superlattice ordering effect in sample C1, suggested by the energy of the InGaP peak at low temperatures, lying at 2.0 eV. [24,25]. Moreover, from the analysis of the low-temperature PCS signal, obtained with higher spectral resolution and focused on the excitation spectral region of quantum-confined states in the QWs, it was possible to separate the light and heavy hole excitons from the continuum: in Fig. 8b, the two contributions are shown in a spectrum recorded at T ¼ 30 K. Here, the hh1e1 excitonic transition appears at higher energy than in the low-temperature PL spectrum, such a significant Stokes shift could be partially induced by the complexity of the interfaces. However, the moderate enlargement of the PCS peak (FWHM ¼ 17 meV, see Fig. 8b) is in favour of a low structural disorder of the MQWs. The hh1e1 and lh1e1 transition energies shift with the temperature by following the GaAs energy gap, as expected (Fig. 8c). Further details of this investigation will be reported elsewhere. Here, we just highlight that, if the hh1e1 and lh1e1 energy transitions are calculated in the limit of T ¼ 0 K, for parabolic bands and by assuming rectangular wells having width of 8 nm and barrier heights consistent with the valence- and conduction-band offsets evaluated in Ref. [17], values about 20 meV higher than the experimental PL peak energies were obtained, after correction for the exciton binding energy (of about 10 meV). This conclusion agrees with PL data obtained previously [14]. Certainly, a slight inadequacy of the above model cannot be excluded, whose possible cause could be the deviation of the wells from the rectangular shape. However, the mentioned discrepancy is significantly reduced if the theoretical expectations are compared with data obtained through excitation-mode measurements, rather than emission measurements. We remark that the absolute energy positions of both the heavy and light hole
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related peaks, obtained by photoelectrical measurements in sample C1, agree with the calculations for a 8.5-nm-wide rectangular QW, in favour of the hypothesis of an effective well width slightly higher than the nominal one, as generally indicated by TEM investigations in similar samples. Despite of the encouraging premises, when a first prototype of device based on the p–i–n structure above discussed was realized, the photovoltaic performances resulted not again satisfactory. The reason of this behaviour is almost certainly related to the background doping of the InGaP alloy in the MQWs. In fact, according to Osborne [26], if the background doping level overcomes the value of 1015 cm3, the electric field profile can significantly vary through the intrinsic region and degrade the performance of the solar cell. Admittance spectroscopy investigation on similar structures revealed that the MQW region is only partially depleted at zero bias. In such conditions, the competition between a short effective diffusion length and the recombination processes probably induce a poor photocurrent value, which implies low conversion efficiency [17]. A reduced thickness of the MQW region could favour a more uniform distribution of the electric field along the structure, but the efficiency would be in this case penalized by a too thin active layer in the device, as can be deduced by the comparison, reported in Ref. [27], of the spectral response obtained from structures containing a different number of QWs. The probable origin of such an unintentional pdoping is the incorporation of carbon during the growth, having already excluded possible effects related to the purity of the carrier gas and of the metalorganic sources. A work is on progress to clarify this point.
5. Conclusions Aiming at the optimum growth conditions, MQW-based devices, formed by p–i–n structure containing 30 MQWs in the intrinsic regions, were realized and studied. The efforts to improve the interface sharpness and to control the residual strain were described in detail in the work and supported by several experimental results. The reproducibility of the satisfactory structural properties and good optical quality of the heterostructures grown at the assessed conditions was undoubtedly shown. A MQW-based test structure was finally fabricated and a spectral photocurrent investigation revealed that it is potentially suitable to work as a photovoltaic cell. Moreover, a well width slightly larger than the nominal value was obtained from the splitting of the PCS emission peaks related to hh1e1 and lh1e1 transitions, consistently with TEM investigation on reference structures. Limitations to solar cell applications of such junctions emerged in terms of performances and are probably related to a still high background doping level of the intrinsic InGaP region.
Acknowledgments
Fig. 8. (a) Room-temperature photocurrent spectrum of the C1 structure, normalized to the response of the experimental setup. (b) Excitonic PC spectrum recorded at T ¼ 30 K with higher resolution than in (a): the hh1e1 lh1e1 excitonic transitions are evidenced by arrows: the peaks appear at E(hh1e1) ¼ 1.543 eV, E(lh1e1) ¼ 1.567 eV, respectively, with a full-width at half-maximum of 17 meV. (c) Temperature dependence of the hh1e1 lh1e1 peak energies (symbols) and of the GaAs energy gap (solid line), the dotted lines are only a guide for the eyes.
This work has been partially supported by the Research Contract EDISON S.p.A./University of Parma, and has been undertaken in the frame of the FIRB/MIUR2007/09 Project ‘‘FOTOENERGIA’’, coordinated by ENEA-Portici (NA). The Authors would like to acknowledge Prof. C. Ghezzi of the Physics Department of the University of Parma for the useful discussions, Dr. L. Nasi of the CNR-IMEM Institute of Parma for the participation to TEM and CL measurements, Dr. E. Gombia of the CNR-IMEM Institute of Parma and Dr. R. Campesato of CESI Milano, for their contribution in processing the samples for the photoelectrical investigation and Dr. M. Baldini of the Physics Department of the University of Parma, for his assistance in the photoelectrical measurements.
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Appendix 1. Supporting Information Supplementary data associated with this article can be found in the online version at doi:10.1016/j.jcrysgro.2009.07.015.
References [1] K. Barnham, I. Ballard, J. Barnes, J. Connolly, P. Griffin, B. Kluftinger, J. Nelson, E. Tsui, A. Zachariou, Appl. Surf. Sci. 113/114 (1997) 722. [2] A. Knauer, P. Krispin, V.R. Balakrishnan, M. Weyers, J. Cryst. Growth 248 (2003) 364. [3] T. Nittono, S. Sugitani, F. Hyuga, J. Appl. Phys. 78 (1995) 5387. ´ [4] R. Kudela, M. Kucˇera, J. Nova k, C. Ferrari, C. Pelosi, J. Cryst. Growth 242 (2002) 132. [5] A. Koizumi, Y. Fujiwara, K. Inoue, T. Yoshikane, A. Urakami, Y. Takeda, Appl. Surf. Sci. 216 (2003) 560. [6] T. Nakano, Y. Nakano, Y. Shimogaki, J. Cryst. Growth 221 (2000) 136. [7] T. Nakano, Y. Nakano, Y. Shimogaki, J. Cryst. Growth 296 (2006) 179. [8] W. Stolz, J. Cryst. Growth 209 (2000) 272. [9] G.B. Stringfellow, Mater. Sci. Eng. B-Solid 87 (2001) 97. [10] A. Moto, S. Tanaka, T. Tanabe, S. Takagishi, Sol. Energy Mater. Sol. Cells 66 (2001) 585. [11] X. Tang, B. Zhang, B. Bo, T. Mei, M.K. Chin, J. Cryst. Growth 288 (2006) 23. [12] B. Mayer, J.P. Reithmaier, A. Forchel, J. Cryst. Growth 227–228 (2001) 298.
¨ [13] S. Leu, F. Hohnsdorf, W. Stolz, R. Becker, A. Salzmann, A. Greiling, J. Cryst. Growth 195 (1998) 98. [14] M. Begotti, M. Longo, R. Magnanini, A. Parisini, L. Tarricone, C. Bocchi, F. Germini, L. Lazzarini, L. Nasi, M. Geddo, Appl. Surf. Sci. 222 (2004) 42. [15] M. Longo, R. Magnanini, A. Parisini, L. Tarricone, C. Bocchi, F. Germini, L. Lazzarini, L. Nasi, EW-MOVPE 2007, 12th European Workshop on Metalorganic Vapour Phase Epitaxy, June 3–6, 2007, Bratislava, Slovakia, Workshop Proceedings, p. 239. [16] D.B. Bushnell, T.N.D. Tibbits, K.W.J. Barnham, J.P. Connolly, M. Mazzer, N.J. Ekins-Daukes, J.S. Roberts, G. Hill, R. Airey, J. Appl. Phys. 97 (2005) 124908. [17] E. Gombia, C. Ghezzi, A. Parisini, L. Tarricone, M. Longo, Mater. Sci. Eng. BSolid 147 (2008) 171. [18] C. Ghezzi, R. Magnanini, A. Parisini, L. Tarricone, E. Gombia, M. Longo, Phys. Rev. B 77 (2008) 125317. [19] D. Taupin, Bull. Soc. Fr. Mineral. Cristallogr. 87 (1964) 469. [20] S. Takagi, Acta Crystallogr. 15 (1962) 1311. [21] X.B. Zhang, J.H. Ryou, R.D. Dupuis, G. Walter, N. Holonyak Jr., J. Electron. Mater. 35 (2006) 705. [22] Y.Q. Wang, Z.L. Wang, T. Brown, A. Brown, G. May, Thin Solid Films 397 (2001) 162. [23] P.J. Collins, K. von Klitzing, K. Ploog, Phys. Rev. B 33 (1986) 4378. ´ [24] M. Longo, A. Parisini, L. Tarricone, L. Toni, R. Kudela, Mater. Sci. Eng. B-Solid 86 (2001) 157. ´ ¨ [25] J. Novak, S. Hasenohrl, R. Kudela, M. Kucera, M.I. Alonso, M. Garriga, Mater. Sci. Eng. B-Solid 88 (2002) 139. [26] J. Osborne, InGaP/GaAs Quantum Well Solar Cells, Imperial College of Science, Technology and Medicine, London, UK, 1994 (M. Sc. Thesis). [27] R. Magnanini, L. Tarricone, A. Parisini, M. Longo, E. Gombia, Thin Solid Films 516 (2008) 6734.
Contents lists available at ScienceDirect
Journal of Crystal Growth journal homepage: www.elsevier.com/locate/jcrysgro
Structural and optical characterization of MOVPE grown InGaP/GaAs MQWs for advanced photovoltaic devices M. Longo a,1, A. Parisini b,, L. Tarricone b, S. Vantaggio b, C. Bocchi c, F. Germini c, L. Lazzarini c a
CNR-INFM, Dipartimento di Fisica, Universita di Parma, Viale G.P. Usberti, 7/A, 43100 Parma, Italy CNISM, Dipartimento di Fisica, Universita di Parma, Viale G.P. Usberti, 7/A, 43100 Parma, Italy c CNR-IMEM Parco Area delle Scienze 37/A, 43010 Loc. Fontanini-Parma, Italy b
a r t i c l e in fo
abstract
Article history: Received 30 January 2009 Received in revised form 12 June 2009 Accepted 5 July 2009 Communicated by K.W. Benz Available online 16 July 2009
The optimization of the electronic properties of InGaP/GaAs MQWs, to be inserted in multilayers heterostructure for novel photovoltaic devices, was performed by structural, optical and photoelectrical measurements. Different sequences of nominally undoped InGaP and GaAs alternated layers were grown by low-pressure metalorganic vapour phase epitaxy, employing tertiarybutylarsine and tertiarybutylphosphine as metalorganic precursors for the V-group elements. In order to minimize the As/P exchange effect, the interface In segregation, and to control the whole lattice matching, single and multi-quantum wells (MQWs) with different: (i) periods, (ii) well widths, (iii) growth temperatures, (iv) gas-switching sequences at the interfaces and (v) indium concentrations in the InGaP alloy, were prepared and investigated. The interface sharpness and the compositional fluctuation of thick MQW region containing up to 40 well-barrier sequences were investigated for the modelling, realization and evaluation of test structures based on low-dimensional systems for third generation solar cells. & 2009 Elsevier B.V. All rights reserved.
PACS: 81.05.Ea 81.15.Gh 81.07.St Keywords: A1. InGaP/GaAs interfaces A3. Gas-switching sequence A3. MOVPE A3. Multi-quantum wells B3. Solar cells B2. Semiconducting indium gallium phosphide
1. Introduction The lattice-matched InGaP/GaAs heterostructure is generally recognized as an interesting system for microelectronic and optoelectronic applications. In particular, it can be inserted in a multilayer structure to obtain a better matching of the solar spectrum, and then enhance the performances of photovoltaic devices. The quantum well solar cells (QWSC) technology, in fact, appears to be a potential candidate to achieve high conversion efficiency devices [1]. The crucial technological requirements for the modelling, realization and test of prototypal photovoltaic devices based on low-dimensional systems, include: (i) as sharp as possible superlattice interfaces, (ii) control of residual lattice strain and structural disorder, (iii) high regularity of the QW periodicity and good correspondence between modelling and Corresponding author. Tel.: +39 0521 905272; fax: +39 0521 905223.
E-mail address: [email protected] (A. Parisini). Present address: Laboratorio Nazionale MDM, Via C. Olivetti, 2, 20041 Agrate Brianza (MI), Italy. 1
0022-0248/$ - see front matter & 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2009.07.015
experimental determination of the bandgap profile, (iv) reduction of background impurity and deep level density. Thus, in the present work, an assessment of the MOVPE growth conditions was carried out to realize QW-based test structures satisfying the (i)–(iv) requirements, to evaluate their potential performances as quantum photovoltaic devices. The advantages of using InGaP with respect to AlGaAs, beyond the possibility to realize Al-free devices, are mainly related to: (i) lower oxidation tendency, (ii) higher mechanical stability and (iii) lower interfacial recombination velocity. On the other hand, when grown by metalorganic vapour phase epitaxy (MOVPE), the InGaP/ GaAs system suffers some problems at the interfaces, such as In memory effects or As/P intermixing. In particular, Indium memory effects and surface segregation can lead to the formation of unintentional InxGa1xAsyP1y layers at the interfaces. Such uncontrolled quaternary interlayer, evidenced by a low-energy undesired optical emission, can affect the lattice match composition and the interfacial morphology, resulting in a degradation of the device properties. Various attempts to improve the interface sharpness of the InGaP/GaAs heterojunction, by exploiting
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different gas switch sequences (GSSs) or by interposing proper interlayers (ILs), have been previously carried out [2–7], the argument being again object of debated research. With regard to safety and purity growth issues, one interesting option is the use of alternative precursors tertiarybutylarsine (TBAs) and tertiarybutylphosphine (TBP), in which one of arsine (phosphine) hydrogen functions is replaced by a heavier butylgroup. In such a way both the toxicity, the oxygen incorporation and the background doping should be significantly lowered by virtue of TBAs and TBP low dissociation temperatures, which also allow the growth below Tg ¼ 600 1C [8,9]. Indeed, (InGaAsP)-based devices, grown by using TBAs and TBP and showing at least comparable otherwise better performances, have been recently realized [10–13]. By employing a MOVPE reactor using TBAs and TBP precursors, we recently investigated the effects of different gas-switching sequences for the growth of multi-quantum wells (MQWs) based on the lattice-matched system InGaP/GaAs [14]. It resulted that the insertion of a few-monolayer thick GaAsP IL at the direct GaAs-on-InGaP interface and the use of proper growth interruption purges lead to optimum optical properties of the structures, by suppressing the undesired anomalous low-energy emission due to In memory effect. Moreover, the emission energies due to transitions between confined quantum states, measured as a function of the well thickness, were in good agreement with the theoretical expectation. In this work, the effects due to the IL insertion and the assessment of growth conditions suitable for the control of the residual strain in the structure are accurately discussed, by taking into account previous results reported in Refs. [14,15] to give a complete view of the work. Optical and photoelectrical investigation on p+-MQW-n test structure are reported, to demonstrate the feasibility of MQW structures with a consistent number of periods, as required for advanced photovoltaic applications [16].
2. Experimentals The growth of the InGaP/GaAs heterostructures was performed by a horizontal low-pressure MOVPE Aixtron reactor (AIX200 RD). In order to optimize the interfaces, previously found conditions were applied [14]. Trimethylgallium (TMGa), trimethylindium, TBAs and TBP metalorganic precursors were used for III–V-group
elements, while hydrogen-diluted disilane was used for n-doping and dimethylzinc for p-doping. Palladium purified H2 was used as a carrier gas. The substrates were 2 inch GaAs (0 0 1) exact wafers, pre-treated by standard etching procedures. The TMGa partial pressure was fixed at 5.6 103 mbar and the V/III ratio was 25.5 for GaAs and 14 for InGaP, the growth was performed at a temperature of 560–600 1C, a total pressure of 50 mbar and a total gas flow of 6.8 l/min. The typical growth rate turned out to be of 1.2 mm/h for GaAs and 2.2 mm/h for InGaP. A particular gasswitching sequence was employed to optimize the optical and morphological quality of the GaAs–InGaP interfaces, as in Ref. [14]. Different sequences of unintentionally doped and nominally pseudomorphic InGaP/GaAs MQWs were grown at Tg ¼ 600 1C, confined within InGaP cladding layers. Single QWs, as well as MQWs, with different periods, variable well size, gas-switching sequence at the interfaces and indium concentration in the InGaP alloy were prepared. In all cases a 200–300-nm-thick GaAs buffer layer was grown on top of the substrate. GaAs wells ranging in thickness between 3 and 10 nm and 80–150 nm thick InGaP barriers were then deposited. The main growth parameters of the samples are collected in Table 1. A background p-type doping of the InGaP layers resulted by electrical investigation of the samples [17,18]. To allow an easier correlation between the results of the present work and those reported in our previous publications for samples featured by same structure and deposition conditions, we point out that the new simple labels here used substitute the identification labels of past publications as follows: samples 1, 2, y stand for samples IGPMQW1, IGPMQW2, y; C1 stands for sample IGPCELL1. The grown structures were investigated through structural and optical measurements. The structural analysis was performed by X-ray diffraction (XRD) and electron microscopy methods. A Philips high-resolution X-ray diffractometer, equipped with a graded multilayer mirror for obtaining a quasi-monochromatic and -parallel beam of high intensity, and a Bartels channel-cut monochromator for selecting the CuKa1 radiation, was used for XRD measurements. To reduce the diffuse scattering coming from the sample, a thin slit was placed in front of the detector with an acceptance angle nearly 0.51 and the diffraction profiles were recorded in the o2y scan mode, by using the 004 Bragg reflection. This arrangement extended the range of the measured
Table 1 Main features of the grown InGaP/GaAs heterostructures. MQW
GaAs well width Nominal (nm)
Measured by TEM (nm)
13
8 8 2 4 6 4 4 4 8
8.370.5 – 3.570.5 5.570.5 7.270.5 5.770.5 4.970.5 4.670.5 –
14
8
–
17 18
8 2 (top) 4 6 8 (bottom) 8
8.670.5 3.071.0 5.070.5 6.570.5 8.570.5 –
1 2 5
6
C1
(top) (middle) (bottom) (top) (middle) (bottom)
GaAsP interlayer (IL)
QW periods
XRD measured indium molar fraction in InGaP
Absent (no IL) IL at direct interface (4 s) Same IL at direct interfaces (4 s)
20 1 3
0.48 – –
Different ILs at direct interfaces (5-3-1 s)
3
0.47
IL at direct+reduced GIT at indirect interface (2 s) IL at direct (4s)+IL at indirect (4 s) interfaces IL at direct interfaces IL at direct interfaces
1
–
1
–
40 20
0.46 0.445
IL at direct interfaces
30
0.488–0.495 (analysis on the whole grown area)
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3. General assessment of MQW structures The milestone of this work was the assessment of the growth parameters for the realization of structures formed by a sequence of several QWs, according to the fundamental requirements for a low-dimensional system-based device. In this section, we recall–giving new details–the main results previously obtained [14,15] and complete the study by adding new experimental data on the final test structure. The achievement of a good morphological quality and a very low structural and compositional disorder had been already confirmed by the DF-TEM image and the XRD curve reported in Ref. [15], resulting in an excellent reproducibility over tens of MQWs. The investigation concerned a MQW heterostructure (sample 1) grown without any particular gas-switching sequence at the interfaces, therefore, without any intentional GaAsP interlayer. The nominal well and barrier thickness of 8 and 12 nm, respectively, show a satisfying agreement with respect to the values of 8.3 and 9.5 nm, measured by TEM images. In this structure, where 20 InGaP/GaAs periods are stacked, the interfaces exhibit a roughness around a few nm, with a lateral period of about 100 nm, slightly increasing from the bottom to the top of the MQWs. Such a lateral unhomogeneity of the interface is also evidenced by a shoulder in the FFT–PL spectrum reported in Ref. [14] (where, the sample is labelled IGP1). However, a shift towards anomalously low values of the PL emission of the MQW reveals the formation of a quaternary InGaAsP layer at the direct GaAs-on-InGaP interface [14].
1 (004)
0
calculated experimental
-1
-2 log (reflectivity)
intensity up to 106 times the value of the substrate Bragg peak. The experimental XRD curves were analysed by means of a best fit procedure based on the X-ray diffraction dynamical theory for distorted crystals [19,20]. The layer thickness t, the static Debye–Waller factor (DW) accounting for the lattice disorder and the lattice spacing modification Dd/d, were the parameters of the calculations. The best fit was obtained by means of a least-squares method for the difference (w) between the experimental and calculated curves. Transmission electron microscopy (TEM) investigations were carried out on /11 0S cross-section specimens by using the (2 0 0) bright field (BF) and dark field (DF) imaging mode, the latter being sensitive to the chemical composition. In this imaging condition, the width of the GaAs quantum wells was measured, by means of proper image analysis software, and compared with the values obtained by HRXRD. Low-temperature cathodoluminescence (CL) investigations were performed on a Cambridge Stereoscan 360 SEM equipped with an Oxford Instruments mono-CL system, with the specimen observed in plan view. Photoluminescence (PL) spectra were detected through a Bruker-IFS66-FFT spectrometer, using the 488 nm line of a 5 W Ar+ ion laser and the 632.8 nm line of a He–Ne laser as excitation and reference light, respectively. A spectral resolution of 0.1 meV (about 1 cm1) was obtained in the investigated spectral range. Finally, photocurrent spectra were obtained through measurements in the temperature range (30–300 K) on p–i–n like structures, whose intrinsic region was formed by a sequence of InGaP/GaAs QWs. The junctions were illuminated by a halogen lamp (100 W), focused into the entrance of a Spex monochromator (0.5 m focal length), with a spectral resolution better than 2 meV. The spectral responses of the MQW p–i–n diodes were obtained by preparing on the top surface a pattern of mesa structures, 500 mm in diameter, by conventional photolithography, as described in Ref. [18]. The PR signal was detected by conventional lock-in techniques and normalized to the system spectral response by using a calibrated Si photodetector.
4295
-3
-4
-5
-6
-7 -3000
-2000
-1000 0 1000 Δθ (arcsec)
2000
3000
Fig. 1. XRD 004 experimental (dot) and calculated (line) profiles of a single QW structure (sample 2) in which a GaAsP interlayer (IL) was inserted at the direct interface.
The efforts to improve the interface quality were subsequently focused on single QWs, where sharp interfaces were obtained when a GaAsP interlayer (IL) is inserted at the direct GaAs-onInGaP interface [14,15]. This expedient was expected to reduce the diffusion/segregation of In through the interface. Fig. 1 shows the XRD spectrum of sample 2, superimposed to the theoretical fit, by which a 1.7 nm thickness was obtained for the GaAsP interlayer, comparable to the nominal value of 1.5 nm. The fair agreement between experiment and theory gives evidence in favour of the high quality of the structure and the good control of the growth parameters. In addition, the PL emission energy of such structure approaches the expectation for the optical transition between the n ¼ 1 heavy hole subband and the n ¼ 1 electron subband in the wells (hh1e1 transition), once the 2D exciton binding energy is taken into account [14]. The effect of either a variation of the growth interruption time (GIT) in sample 13 or the insertion of a GaAsP IL in sample 14 at the indirect interface were also investigated, without varying the well and barrier thickness. As it is shown in Fig. 2, in both cases the sharpness of the interface degrades with respect to sample 2, in fact, low-energy peaks appears again in the PL spectrum of samples 13 and 14, also in this case attributable to the formation of an undesired quaternary alloy [21]. In addition, the main emission peak, detected, as expected, at about 1.535 eV in samples 13 and 2, shifts toward higher energies (1.557 eV) in sample 14. This peak is due to the hh1e1 excitonic recombination, whose energy position depends on the effective well width, therefore, an apparent reduction of the QW thickness from 8 to 5 nm seems to be induced in the QW of sample 14. 3.1. Optimization of the GaAsP IL thickness at the direct interface With the aim to more deeply investigate the role of the GaAsP IL, two samples containing a different sequence of three stacked
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average over the three wells, showing a single emission peak, with a minor effect on the linewidth. The active role of the IL in limiting the As/P intermixing and the In diffusion was, however, confirmed by XRD analysis on sample 6 (Fig. 4). Also, a good control of the lattice match at the InGaP/GaAs interfaces resulted again, with /xInS ¼ 0.47 for each barrier. No effects, ascribable to the interface broadening were observed, the roughness being in any case again limited within 1 nm, with the upper (indirect) interface
Fig. 2. FFT–PL spectrum recorded at 12 K from MQW structures grown by different gas-switching sequences (GSSs) at the indirect interfaces. Arrows indicate in each spectrum the hh1e1 excitonic transition.
Fig. 3. CL spectra of sample 6, recorded at different probing depths (electron accelerating voltages). The peaks are relative to the hh1e1 excitonic transition.
1 fitting experimental 0
(004)
-1
log (reflectivity)
InGaP/GaAs QWs were grown, that is samples 5 and 6, having either the same GaAsP IL thickness (expected thickness around 1.7 nm). Sample 5 contains three single QWs of different widths (6, 4 and 2 nm, from the bottom to the top, respectively), sample 6 contains three single QWs of nominally identical width (4 nm), but different GaAsP ILs thickness (0.4, 1.2 and 2.0 nm, from the bottom to the top, respectively). In both samples the barriers were 150 nm thick. The well-width values obtained by TEM in sample 5, where the IL was four monolayers thick, were found to be slightly higher than the nominal ones, that is 7.270.5, 5.570.5 and 3.570.5 nm, respectively. The low-temperature PL spectrum of this sample, already shown in Fig. 2b of Ref. [14], exhibits three well-separated peaks, lying at 1.545, 1.571 and 1.620 eV, respectively, which quite well follow the theoretical trend expected for the hh1e1 radiative transition. The slight but systematic red shift of the optical emission data with respect to the theoretical energies, in spite of the inclusion of the 2D exciton binding energy, will be briefly discussed later. The TEM data seem to suggest that the presence of the IL only slightly influences the well width. However, the investigation of sample 6 revealed that, when the IL thickness increases to 2 nm, the morphology of the well degrades and the effective well width increases from a nominal value of 4 nm to an effective value of 5.771 nm, significantly more with respect to the other two quantum wells, which appear 4.970.5 and 4.670.5 nm wide, for a 1.2 and 0.4-nm-thick IL, respectively [14]. This result was also confirmed by space-resolved cathodoluminescence (CL) measurements. By taking advantage of the dependence of the beam penetration depth on the kinetic energy of the exciting electrons, and with the aid of Montecarlo simulations (not shown here), we were able to selectively investigate the different wells through the variation of the voltage accelerating the incident electrons. As it appears from Fig. 3, the 3 kV CL spectrum, obtained for sample 6 from the top well prevalently, shows a main emission peak (E ¼ 1.567 eV) at a slightly lower energy with respect to the energies of the emission peaks dominating the CL spectrum, both at an intermediate accelerating voltage (V ¼ 4 kV, emission from the middle QW) and at the highest accelerating voltage (V ¼ 5 kV, emission from the bottom QW). This behaviour suggests a larger width of the top well with respect to the other two. Nevertheless, such a spread is too limited to be evidenced by PL measurement, which gives an
-2
-3
-4
-5
-6 -2000
-1000
0 Δθ [arcsec]
1000
2000
Fig. 4. Experimental (dot) and calculated (line) 004 X-ray diffraction profiles of sample 6.
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(ii)
2.0 (iii)
PL intensity (a.u.)
T = 10 K 1.5
(iv)
1.0
0.5
0.0 1.2
1.3
1.4
1.5 Energy (eV)
1.6
1.7
(i) Lw = 2 nm
(ii) Lw = 4 nm
(iii) Lw = 6 nm
(iv) Lw = 8 nm
Fig. 5. PL peak related to the hh1e1 excitonic transition at T ¼ 10 K (a) and BF TEM image (b) acquired from a lattice mismatched structure containing 40 GaAs quantum wells, 8 nm thick (sample 17).
Fig. 6. PL spectrum at T ¼ 10 K (a) and (2 0 0) DF TEM image (b) of a lattice mismatched structure containing four series of quantum wells, each containing four periods, having well width decreasing from 8 to 2 nm (sample 18). Only the hh1e1 excitonic PL peaks relative to the MQWs having well width, respectively, of (ii) 4 nm, (iii) 6 nm and (iv) 8 nm are detected. In the image (b), the corresponding sequence of MQWs is evidenced.
of the wells worse than the lower (direct) one, and no lattice defect due to dislocations being evidenced [14,15].
3.2. Lattice mismatch effects The requirement of growing heterostructures containing a consistent number of QW periods imposes an accurate control of the InGaP/GaAs lattice mismatch and the effects of the residual strain. Different MQW structures were grown, having different stacking sequences and well widths. In particular, we now discuss the results relative to a series of 40 QWs, having the same well width Lw ¼ 8 nm (sample 17), and a structure, sample 18, formed by four groups of MQWs (of four QWs each), having Lw decreasing from 8 nm (in the bottom QWs), to 6, 4 and 2 nm (in the top QWs). At each direct interface, a 1.7-nm-thick GaAsP IL was interposed. Fig. 5a and b, respectively, show the low-temperature FFT–PL spectrum and the TEM image obtained for sample 17, whose In mole fraction resulted intentionally low, xIn ¼ 0.46, with respect to the perfect lattice matching value of xIn ¼ 0.484. The main peak in the PL spectrum at E ¼ 1.531 eV, with a FWHM of 0.013 eV, is consistent with the expectations, while the weak shoulder
observed at E ¼ 1.46 eV could be attributed to interface defects or non-intentional impurities. The TEM image shows a strong modification of the MQW periodicity after about 70 interfaces, due to the elastic strain relaxation. Because of the residual lattice mismatch, the structure is locally perturbed but has not relaxed as a whole and misfit dislocations have not been detected at any interface. A similar analysis on sample 18 reveals three PL emission peaks due to the first three sets of QWs (Fig. 6a), at E ¼ 1.533, 1.546 and 1.563 eV from the bottom to the top of QWs set, whereas the 2 nm wide QW is completely destroyed by the lattice relaxation (Fig. 6b), which in this structure appears after 30 interfaces, again because of a low In content (In mole fraction ¼ 0.445). The elastic energy accumulated in the structure partially favours the interface roughness, in fact, in the literature the latter has been considered as the main cause of the degradation of the top layers in InGaP/GaAs MQWs. Indeed, this effect has been evidenced also in lattice-matched structures containing many interfaces [22].
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Table 2 Scheme of the optimized structure C1, containing 30 staked MQWs embedded in doped InGaP cladding layers. Layer
Material
Thickness (nm)
Dopant type
Dopant
Dopant concentration (cm3)
Cap Top cladding layer Top spacer layer Barrier Well Direct interface IL Bottom spacer layer Bottom cladding layer Buffer Substrate
GaAs InGaP InGaP InGaP GaAs GaAsP InGaP InGaP GaAs GaAs
10 300 50 12 8 1.5 100 300 250
Zn Zn
p++ p+
2 1018 1018
Si Si Si
n+ n+ n+
2 1018 2 1018
30 periods
Fig. 7. FFT–PL spectrum at T ¼ 10 K obtained from the optimized p–i–n structure containing a MQW of 30 period (sample C1).
4. Study of p+-MQW-n+ test structures The assessment of the growth parameters was exploited for the design and realization of a p–i–n test structures containing MQWs in the intrinsic region (sample C1), which is the basic structure for the realization of novel photovoltaic devices (QWSCs). Thirty lattice-matched InGaP/GaAs QWs were stacked. In order to limit impurity diffusion from the cladding layers into the MQW region, in this case, spacer layers were properly interposed. The details of the structure are reported in Table 2. A very high optical quality of the sample appears from the low-temperature (12 K) FFT–PL spectrum, dominated by the narrow emission from the Lw ¼ 8 nm width wells (Fig. 7). The PL peak, due to the hh1e1 transition, lies at E ¼ 1.530 eV and has a full-width at half-maximum (FWHM) of 12 meV. The energy peak position was also confirmed by a CL investigation, indicating a good reproducibility over all the samples prepared with the same modelling parameters. The optical properties of the test structures were more extensively investigated through photocurrent spectroscopy (PCS), which gives information complementary to PL and CL. In fact, whereas the latter techniques reveal the more probable optical recombination process (emission of photons), in the PCS, photons of variable wavelength are absorbed, giving information on permitted excitation transitions. The PCS investigation of MQWs, in particular, evidences the effects of spatial confinement
of carriers and offers the possibility to compare nominal and effective QW sizes [23]. In addition, it is strongly influenced by the collection efficiency of carriers at the contacts, then by the transport mechanisms of photo-generated carriers through the QWs. Fig. 8a shows the photocurrent spectrum of sample C1, recorded at T ¼ 300 K. The spectrum, extends between the values of the GaAs and InGaP energy gaps (window effect), as expected. The shoulders observed at low energies are related to the hh1e1 and lh1e1 excitonic transitions; generally, the evidence of the splitting between heavy and light hole related peaks is a fingerprint of an optimum structural quality of the MQW, since a high density of defects and the internal field due to the presence of strain generally broaden the optical features, resulting in an enlarged lineshape. Some additional information on the structural quality of the samples can be obtained by the analysis of the whole PC spectrum at different temperatures. As a minor feature, we mention the evidence of a negligible superlattice ordering effect in sample C1, suggested by the energy of the InGaP peak at low temperatures, lying at 2.0 eV. [24,25]. Moreover, from the analysis of the low-temperature PCS signal, obtained with higher spectral resolution and focused on the excitation spectral region of quantum-confined states in the QWs, it was possible to separate the light and heavy hole excitons from the continuum: in Fig. 8b, the two contributions are shown in a spectrum recorded at T ¼ 30 K. Here, the hh1e1 excitonic transition appears at higher energy than in the low-temperature PL spectrum, such a significant Stokes shift could be partially induced by the complexity of the interfaces. However, the moderate enlargement of the PCS peak (FWHM ¼ 17 meV, see Fig. 8b) is in favour of a low structural disorder of the MQWs. The hh1e1 and lh1e1 transition energies shift with the temperature by following the GaAs energy gap, as expected (Fig. 8c). Further details of this investigation will be reported elsewhere. Here, we just highlight that, if the hh1e1 and lh1e1 energy transitions are calculated in the limit of T ¼ 0 K, for parabolic bands and by assuming rectangular wells having width of 8 nm and barrier heights consistent with the valence- and conduction-band offsets evaluated in Ref. [17], values about 20 meV higher than the experimental PL peak energies were obtained, after correction for the exciton binding energy (of about 10 meV). This conclusion agrees with PL data obtained previously [14]. Certainly, a slight inadequacy of the above model cannot be excluded, whose possible cause could be the deviation of the wells from the rectangular shape. However, the mentioned discrepancy is significantly reduced if the theoretical expectations are compared with data obtained through excitation-mode measurements, rather than emission measurements. We remark that the absolute energy positions of both the heavy and light hole
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related peaks, obtained by photoelectrical measurements in sample C1, agree with the calculations for a 8.5-nm-wide rectangular QW, in favour of the hypothesis of an effective well width slightly higher than the nominal one, as generally indicated by TEM investigations in similar samples. Despite of the encouraging premises, when a first prototype of device based on the p–i–n structure above discussed was realized, the photovoltaic performances resulted not again satisfactory. The reason of this behaviour is almost certainly related to the background doping of the InGaP alloy in the MQWs. In fact, according to Osborne [26], if the background doping level overcomes the value of 1015 cm3, the electric field profile can significantly vary through the intrinsic region and degrade the performance of the solar cell. Admittance spectroscopy investigation on similar structures revealed that the MQW region is only partially depleted at zero bias. In such conditions, the competition between a short effective diffusion length and the recombination processes probably induce a poor photocurrent value, which implies low conversion efficiency [17]. A reduced thickness of the MQW region could favour a more uniform distribution of the electric field along the structure, but the efficiency would be in this case penalized by a too thin active layer in the device, as can be deduced by the comparison, reported in Ref. [27], of the spectral response obtained from structures containing a different number of QWs. The probable origin of such an unintentional pdoping is the incorporation of carbon during the growth, having already excluded possible effects related to the purity of the carrier gas and of the metalorganic sources. A work is on progress to clarify this point.
5. Conclusions Aiming at the optimum growth conditions, MQW-based devices, formed by p–i–n structure containing 30 MQWs in the intrinsic regions, were realized and studied. The efforts to improve the interface sharpness and to control the residual strain were described in detail in the work and supported by several experimental results. The reproducibility of the satisfactory structural properties and good optical quality of the heterostructures grown at the assessed conditions was undoubtedly shown. A MQW-based test structure was finally fabricated and a spectral photocurrent investigation revealed that it is potentially suitable to work as a photovoltaic cell. Moreover, a well width slightly larger than the nominal value was obtained from the splitting of the PCS emission peaks related to hh1e1 and lh1e1 transitions, consistently with TEM investigation on reference structures. Limitations to solar cell applications of such junctions emerged in terms of performances and are probably related to a still high background doping level of the intrinsic InGaP region.
Acknowledgments
Fig. 8. (a) Room-temperature photocurrent spectrum of the C1 structure, normalized to the response of the experimental setup. (b) Excitonic PC spectrum recorded at T ¼ 30 K with higher resolution than in (a): the hh1e1 lh1e1 excitonic transitions are evidenced by arrows: the peaks appear at E(hh1e1) ¼ 1.543 eV, E(lh1e1) ¼ 1.567 eV, respectively, with a full-width at half-maximum of 17 meV. (c) Temperature dependence of the hh1e1 lh1e1 peak energies (symbols) and of the GaAs energy gap (solid line), the dotted lines are only a guide for the eyes.
This work has been partially supported by the Research Contract EDISON S.p.A./University of Parma, and has been undertaken in the frame of the FIRB/MIUR2007/09 Project ‘‘FOTOENERGIA’’, coordinated by ENEA-Portici (NA). The Authors would like to acknowledge Prof. C. Ghezzi of the Physics Department of the University of Parma for the useful discussions, Dr. L. Nasi of the CNR-IMEM Institute of Parma for the participation to TEM and CL measurements, Dr. E. Gombia of the CNR-IMEM Institute of Parma and Dr. R. Campesato of CESI Milano, for their contribution in processing the samples for the photoelectrical investigation and Dr. M. Baldini of the Physics Department of the University of Parma, for his assistance in the photoelectrical measurements.
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Appendix 1. Supporting Information Supplementary data associated with this article can be found in the online version at doi:10.1016/j.jcrysgro.2009.07.015.
References [1] K. Barnham, I. Ballard, J. Barnes, J. Connolly, P. Griffin, B. Kluftinger, J. Nelson, E. Tsui, A. Zachariou, Appl. Surf. Sci. 113/114 (1997) 722. [2] A. Knauer, P. Krispin, V.R. Balakrishnan, M. Weyers, J. Cryst. Growth 248 (2003) 364. [3] T. Nittono, S. Sugitani, F. Hyuga, J. Appl. Phys. 78 (1995) 5387. ´ [4] R. Kudela, M. Kucˇera, J. Nova k, C. Ferrari, C. Pelosi, J. Cryst. Growth 242 (2002) 132. [5] A. Koizumi, Y. Fujiwara, K. Inoue, T. Yoshikane, A. Urakami, Y. Takeda, Appl. Surf. Sci. 216 (2003) 560. [6] T. Nakano, Y. Nakano, Y. Shimogaki, J. Cryst. Growth 221 (2000) 136. [7] T. Nakano, Y. Nakano, Y. Shimogaki, J. Cryst. Growth 296 (2006) 179. [8] W. Stolz, J. Cryst. Growth 209 (2000) 272. [9] G.B. Stringfellow, Mater. Sci. Eng. B-Solid 87 (2001) 97. [10] A. Moto, S. Tanaka, T. Tanabe, S. Takagishi, Sol. Energy Mater. Sol. Cells 66 (2001) 585. [11] X. Tang, B. Zhang, B. Bo, T. Mei, M.K. Chin, J. Cryst. Growth 288 (2006) 23. [12] B. Mayer, J.P. Reithmaier, A. Forchel, J. Cryst. Growth 227–228 (2001) 298.
¨ [13] S. Leu, F. Hohnsdorf, W. Stolz, R. Becker, A. Salzmann, A. Greiling, J. Cryst. Growth 195 (1998) 98. [14] M. Begotti, M. Longo, R. Magnanini, A. Parisini, L. Tarricone, C. Bocchi, F. Germini, L. Lazzarini, L. Nasi, M. Geddo, Appl. Surf. Sci. 222 (2004) 42. [15] M. Longo, R. Magnanini, A. Parisini, L. Tarricone, C. Bocchi, F. Germini, L. Lazzarini, L. Nasi, EW-MOVPE 2007, 12th European Workshop on Metalorganic Vapour Phase Epitaxy, June 3–6, 2007, Bratislava, Slovakia, Workshop Proceedings, p. 239. [16] D.B. Bushnell, T.N.D. Tibbits, K.W.J. Barnham, J.P. Connolly, M. Mazzer, N.J. Ekins-Daukes, J.S. Roberts, G. Hill, R. Airey, J. Appl. Phys. 97 (2005) 124908. [17] E. Gombia, C. Ghezzi, A. Parisini, L. Tarricone, M. Longo, Mater. Sci. Eng. BSolid 147 (2008) 171. [18] C. Ghezzi, R. Magnanini, A. Parisini, L. Tarricone, E. Gombia, M. Longo, Phys. Rev. B 77 (2008) 125317. [19] D. Taupin, Bull. Soc. Fr. Mineral. Cristallogr. 87 (1964) 469. [20] S. Takagi, Acta Crystallogr. 15 (1962) 1311. [21] X.B. Zhang, J.H. Ryou, R.D. Dupuis, G. Walter, N. Holonyak Jr., J. Electron. Mater. 35 (2006) 705. [22] Y.Q. Wang, Z.L. Wang, T. Brown, A. Brown, G. May, Thin Solid Films 397 (2001) 162. [23] P.J. Collins, K. von Klitzing, K. Ploog, Phys. Rev. B 33 (1986) 4378. ´ [24] M. Longo, A. Parisini, L. Tarricone, L. Toni, R. Kudela, Mater. Sci. Eng. B-Solid 86 (2001) 157. ´ ¨ [25] J. Novak, S. Hasenohrl, R. Kudela, M. Kucera, M.I. Alonso, M. Garriga, Mater. Sci. Eng. B-Solid 88 (2002) 139. [26] J. Osborne, InGaP/GaAs Quantum Well Solar Cells, Imperial College of Science, Technology and Medicine, London, UK, 1994 (M. Sc. Thesis). [27] R. Magnanini, L. Tarricone, A. Parisini, M. Longo, E. Gombia, Thin Solid Films 516 (2008) 6734.