Semicond. Sci. Technol. 13 (1998) 936–940. Printed in the UK

PII: S0268-1242(98)92290-8

Long-wavelength (Ga, In)Sb/GaSb strained quantum well lasers grown by molecular beam epitaxy N Bertru†‡, A Baranov†, Y Cuminal†, G Almuneau†, F Genty†, ´ O Brandt§, A Mazuelas§ and K H Ploog§ A Joullie†, † CEM2, Case Courier 067, Universite´ de Montpellier II, UMR CNRS n◦ 5507, ` Place Eugene Bataillon, F-34095 Montpellier Cedex 05, France ¨ § Paul-Drude-Institut fur Hausvogteiplatz 5-7, D-10117 Berlin, ¨ Festkorperelektronik, Germany Received 6 March 1998, accepted for publication 27 April 1998 Abstract. The molecular beam epitaxy growth of strained (Ga, In)Sb/GaSb quantum wells is investigated. In a narrow range of growth conditions, (Ga, In)Sb quantum well structures exhibiting excellent structural properties as well as intense and narrow photoluminescence transitions are obtained. Stimulated emission at 1.98 µm is observed at room temperature from laser diodes with Ga0.74 In0.26 Sb/GaSb strained quantum wells as the active zone. The lasers exhibit threshold current densities as low as 280 A cm−2 and a characteristic temperature of 75 K.

1. Introduction High-quality laser diodes emitting near 2 µm at room temperature (RT) can be used in a wide range of applications including chemical sensing for atmospheric pollution and drug monitoring, ‘eye safe’ communications and medical procedures such as laser surgery and medical diagnostics. Another important application is optical pumping of Ho and Ho:Th doped YAG solid state lasers emitting at 2.1 µm. High-power RT laser diodes operating in the 2 µm wavelength range have been fabricated from strained multi-quantum-well (MQW) laser structures made from either the (Ga, In)As/(Ga, In)(As, P) [1–3] or the (Ga, In)(As, Sb)/(Al, Ga)(As, Sb) [4–6] material system. More recently, low-threshold high-performance (Ga, In)(As, Sb)/GaSb MQW laser diodes emitting above 2 µm have been reported [7]. The use of Sb-based lasers is promising because of the possibility to extend the operating wavelength range well above 2 µm. Nevertheless, the growth of Sb-based quantum structures by molecular beam epitaxy (MBE) has not been studied in detail yet. Very little work has been reported on the growth of (Ga, In)Sb/GaSb strained quantum wells which can be considered as a model system for studying strain effects in Sb compounds [8–15]. In this paper, we investigate the MBE growth and the optical properties of (Ga, In)Sb/GaSb compressively strained QWs. First, we report reflection high-energy ‡ Present address: Laboratoire de Physique du Solide (INSA), Avenue des Buttes de Coesmes F-35043 Rennes Cedex, France. c 1998 IOP Publishing Ltd 0268-1242/98/080936+05$19.50

electron diffraction (RHEED) oscillation studies for accurate calibration of the Sb flux. Next, the effect of the growth temperature on the crystalline and optical properties of (Ga, In)Sb/GaSb strained QWs are studied. We show that in a narrow range of growth temperatures, both high optical and structural quality can be achieved. Finally, we demonstrate the possibility of achieving lasing at room temperature from laser structures including (Ga, In)Sb/GaSb MQWs in the active zone. 2. Experiments Growth runs have been performed on two different MBE systems. The samples for structural and optical investigation were grown in a Riber system equipped with solid sources. The laser structures were grown in a second solid-source MBE system (Varian), equipped with Sb2 Te3 and Be cells for n- and p-type doping respectively. The substrate temperatures were measured by a thermocouple contacting the backside of the holders. Reproducible temperature measurements in both MBE systems were achieved by observing the (1 × 3) to (1 × 5) reconstruction transition on GaSb which occurs at 400 ◦ C under an Sb effective flux of 0.8 monolayer per second (ML s−1 ). The x-ray diffraction experiments were performed using a double-crystal diffractometer with a Cu anode ˚ and a (100) asymmetrically [l(Cu Kα1 ) = 1.540 56 A] cut Ge crystal as monochromator and collimator. An HeNe laser was used as an excitation source for the

MBE-grown LW (Ga, In)Sb/GaSb strained QW lasers

Figure 1. RHEED intensity recorded during growth of (a) GaSb and (b) (Ga, In)Sb under Sb-rich conditions and (c) Sb-induced intensity RHEED oscillations at 400 ◦ C. The inset shows the effective incorporation rate of Sb deduced from RHEED oscillations (full circles) and from the AlSb reconstruction transition observed at 550 ◦ C (open circles).

Figure 2. Experimental (dots) and simulated (line) XRD profiles of Ga0.9 In0.1 Sb/GaSb compressively strained QWs grown at 425 ◦ C.

photoluminescence (PL) measurements. The luminescence signal was detected with a liquid nitrogen cooled InSb detector.

3. RHEED studies of the MBE growth of GaSb and (Ga, In)Sb The RHEED oscillation technique has proved to be a very powerful technique for studying and controlling MBE growth. However, in the case of antimonides the use

of this technique is not yet well developed due to the difficulties of observing well defined RHEED oscillations. The large surface diffusion length of physisorbed In and Ga on the GaSb surface, which favours the step flow mode, is the origin of these difficulties [16, 17]. RHEED intensity oscillations which we recorded during the growth of GaSb and (Ga, In)Sb at 400 ◦ C are shown in figure 1, curves (a) and (b). Many oscillations of large amplitude are observed. We assume that the comparatively straightforward detection of RHEED intensity oscillations in this work is partly due to our low growth temperatures, but above all to the UV 937

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O3 treatment [18] of the GaSb substrates. Indeed, this technique permits us to prepare surfaces with lower step densities than the usual wet etching method [19], and a significant improvement of RHEED intensity oscillations is noticed on samples prepared by this technique. Stoichiometric growth of InSb and GaSb cannot occur when a large excess of group V elements is supplied to the surface, contrary to the case of arsenides. In fact, the excess of Sb on the surface enhances the formation of defects, such as vacancies and clusters, and has a detrimental effect on the optical properties of the samples [20]. Thus, an accurate control of Sb flux (fSb ) is crucial to achieve highquality antimonide layers. Sb-induced RHEED intensity oscillations recorded at 400 ◦ C are shown in figure 1 curve (c). The excess of Ga on the surface was accomplished by establishing fSb /fGa < 1. Under these conditions, the period of the first oscillations is not constant, since the surface gradually changes from Sb to Ga stabilization. Then, a steady state (constant period) is reached for which each oscillation corresponds to the incorporation of one monolayer of Sb. The effective incorporation rates measured for various cell temperatures are presented in the inset of figure 1. An Arrhenius behaviour of the Sb flux versus cell temperature is observed as expected for a conventional Knudsen source. Another procedure for the estimation of the Sb flux has been proposed by Subanna et al [21]. This method is based on the transition of a (1 × 3) Sb-stable to a (4 × 2) Al-stable reconstruction on AlSb, which occurs, according to these authors, at equal Al and Sb fluxes at 550 ◦ C. We employed this procedure via a gradual increase of the Al flux during the growth of AlSb, assuming that the effective Sb incorporation rate equals the AlSb growth rate when the surface reconstruction transition occurs. The Sb fluxes measured in this manner are shown in the inset of figure 1. Excellent agreement with the values deduced from Sb-induced RHEED intensity oscillations is obtained. Subanna’s procedure thus seems to be a very useful technique for fast and accurate measurement of the Sb flux. Moreover, the absence of Sb desorption at high temperatures (550 ◦ C) shows that the surface residence time of Sb is long compared with the case of As. This experimental finding is important for the understanding of the incorporation of Sb which represents a crucial aspect of antimonide growth [22]. 4. Optimum growth temperatures of (Ga, In)Sb/GaSb QWs In order to determine the optimum growth temperature of (Ga, In)Sb/GaSb strained QWs, a set of samples has been grown on (100) GaSb substrates at temperatures ranging from 300 to 550 ◦ C. Their structure consisted of 6 QWs ˚ thick GaSb barrier layers. The well separated by 350 A ˚ and 10% thickness and In mole fraction were set to 60 A respectively. These values are far below those expected for the formation of misfit dislocations, and were chosen to avoid interpretation problems related to plastic relaxation. A striking peculiarity of the x-ray diffraction (XRD) profiles of these samples, as shown in figure 2, is the existence 938

Figure 3. PL spectrum taken at 5 K from the Ga0.9 In0.1 Sb/GaSb QW sample grown at 400 ◦ C with an Sb/Ga flux ratio of 1.7.

Figure 4. Evolution of the integrated luminescence intensity and of the PL peak linewidth recorded at 5 K with growth temperature. The curves are a guide to the eye.

of two peaks corresponding to the GaSb substrate and the GaSb buffer layer. We attribute the observed mismatch to the presence of As in the GaSb buffer layers. From Vegard’s law, assuming a completely strained GaSb buffer layer on the GaSb substrate, we calculate an As mole fraction of less than 0.005 for all the samples. The origin of this As contamination is probably related to the previous growth of arsenides in this MBE system. In addition to the peaks related to the GaSb substrate and the GaSb buffer layer, the XRD profiles of samples grown at lower temperatures (450 ◦ C) satellite peaks up to the fifth order. This observation, together with the presence of clearly resolved pendell¨osung fringes, demonstrates that the interfaces are smooth and that the pseudomorphic alloy layers are of high crystalline quality for the samples grown at temperatures lower than 450 ◦ C. Simulated profiles which were computed taking into account the As contamination show excellent agreement with the experimental profiles (figure 2). This fact demonstrates the high periodicity of the (Ga, In)Sb/GaSb QW structures. A PL spectrum recorded at 5 K from a sample grown at 400 ◦ C is shown in figure 3. Three different transitions are observed, from which the two peaks highest in energy

MBE-grown LW (Ga, In)Sb/GaSb strained QW lasers

Figure 5. Schematic energy band diagram for a (Ga, In)Sb laser structure emitting at 2 µm.

(at 0.778 eV and 0.79 eV) are known to be due to a conduction band–native acceptor and an as yet unidentified transition in GaSb respectively [23]. The dominating narrow (≈10 meV) peak at 0.750 eV originates from the (Ga, In)Sb QWs. The ratio between this transition and those related to the GaSb barriers is greater than 40, which constitutes a significant improvement compared with previously reported results [12, 13]. The linewidth and the integrated intensity of the QW line for samples grown between 300 and 550 ◦ C are shown in figure 4. The samples grown at 350 ◦ C and at lower temperatures exhibit a very low PL efficiency. In this low temperature range, Bata et al [24] have observed Sb condensation during the growth of InSb, and the inferior optical properties of these samples may thus be related to the formation of Sb clusters. This interpretation is partially confirmed by the abrupt enhancement of the PL efficiency on samples grown around 400 ◦ C. The narrowest linewidths and strongest PL intensities are obtained for samples grown in this temperature range. When further increasing Ts , a decrease of the PL efficiency is observed. This effect is presumabely related to the low melting point of antimonides (TGaSb = 712 ◦ C, TInSb = 524 ◦ C) which favours defect formation at relatively low temperatures [25]. The relatively low growth temperature used in MBE seems to be an advantage in comparison with metal-organic chemical vapour deposition (MOCVD) which requires a high temperature (600 ◦ C) in order to achieve complete decomposition of the precursors [26]. 5. Laser fabrication and characteristics In order to test the quality of the (Ga, In)Sb/GaSb QWs, we prepared the laser structure shown schematically in figure 5 together with a corresponding band edge diagram. The following epitaxial layers were grown on (100)n-GaSb substrates: 0.2 µm thick graded n-AlGaAsSb buffer layer, 2 µm thick n-Al0.55 Ga0.45 As0.04 Sb0.96 cladding, 0.6 µm thick nominally undoped active zone (see below), 2 µm thick Al0.55 Ga0.45 As0.04 Sb0.96 p-cladding, and 0.07 µm thick p+ -GaSb cap. Since AlSb and GaSb have a lattice mismatch of 6.5 × 10−3 (about four times larger than that

Figure 6. Threshold current as a function of temperature. The inset shows the lasing spectrum at RT for an injection current I ≈ 1.2Ith . The line is an exponential fit.

between AlAs and GaAs), a small amount of As (≈4%) was intentionally incorporated in the Al0.55 Ga0.45 As0.04 Sb0.96 cladding layers to obtain lattice matching with the substrate and to avoid formation of misfit dislocation. The active zone consisted of four compressively strained (1a/a ≈ 1.6%) Ga0.74 In0.26 Sb/GaSb quantum wells. The thickness of the QWs and barrier layers was set to 60 and 300 nm respectively. During the growth, the substrate temperature was set at 530 ◦ C for the Al0.55 Ga0.45 As0.04 Sb0.96 layers and at 400 ◦ C for the QWs, the latter temperature having been previously identified as the optimum growth temperature. Broad area (≈200 µm wide) Fabry–P´erot laser diodes were fabricated from the grown structures. The cleaved facets were uncoated and the cavity length was around 500 µm. The devices were tested under pulsed excitation at temperatures close to 300 K. RT laser emission with a few longitudinal modes centred at 1.98 µm was observed (inset of figure 6) under quasi cw injection (40% duty cycle). The threshold current density was found to be 280 A cm−2 at RT and increased with temperature between 240 K and 350 K as exp(T /T0 ) with a characteristic temperature T0 of 75 K, as shown in figure 6. This relatively low T0 value can be attributed to a poor hole confinement by the GaSb barriers layers as 939

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also observed by PL spectroscopy on (Ga, In)Sb/GaSb QWs [27]. In order to improve the hole confinement, (Al, Ga)Sb alloys may be used as the barrier material. However, it is well known that high-quality Al-containing compounds are difficult to prepare. Indeed, the threshold current density of our (Ga, In)Sb/GaSb lasers is comparable with the best values from Al-containing lasers emitting near 2 µm [4–6]. Further studies have to be performed to determine the optimal compromise between good hole confinement and high material quality. 6. Conclusions Strained (Ga, In)Sb/GaSb QWs have been grown on GaSb substrates by MBE. Sb-induced oscillations and the AlSb reconstruction transition were used to accurately determine the Sb flux. The excellent agreement observed shows that the latter pocedure based on the AlSb reconstruction transition is a useful method for estimating the Sb flux. We have shown that growth in a narrow range of temperatures (around 400 ◦ C) results in the achievement of QWs which exhibit excellent structural properties as well as intense and narrow photoluminescence peaks. Finally, the (Ga, In)Sb/GaSb QWs were inserted in a laser structure. Stimulated emission at 1.98 µm at room temperature was observed for lasers with four Ga0.74 In0.26 Sb/GaSb strained QWs in the active zone. The lasers exhibit threshold current densities around 280 A cm−2 and a characteristic temperature of 75 K. Acknowledgments We would like to thank G Paris for the PL measurements, D Spaniol for the chemical preparation of the GaSb substrates and J L Lazzari for fruitful discussions. References [1] Forouhar S, Ksendov A, Larsson A G and Temkin H 1992 Electron. Lett. 28 1431 [2] Major J S, Nam D W, Osinski J S and Welsh D F 1993 IEEE Photon. Technol. Lett. 5 594

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[3] O’Brien S, Plano W, Major J, Welch D F and Tally T 1995 Electron. Lett. 31 105 [4] Choi H K and Eglash S J 1992 Appl. Phys. Lett. 61 1154 [5] Choi H K, Turner G W and Eglash S J 1994 IEEE Photon. Technol. Lett. 6 7 [6] Garbuzov D Z, Martinelli R U, Lee H, Menna R J, York P K, Di Marco L A, Harvey M G, Matarese R J, Narayan S Y and Connolly J C 1997 Appl. Phys. Lett. 70 2931 [7] Baranov A N, Cuminal Y, Boissier G, Alibert C and Joulli´e A 1996 Electron. Lett. 32 2279 [8] Kodama M 1994 Phys. Status Solidi a 141 145 [9] Chen S M, Su Y K and Lu Y T 1993 J. Appl. Phys. 74 7288 [10] Jiang Y, Teich M C and Wang W I 1990 Appl. Phys. Lett. 57 2922 [11] Haywood S K, Chidley E T R, Mallard R E, Mason N J, Nicholas R J, Walker P J and Warburton R J 1989 Appl. Phys. Lett. 54 922 [12] Qian L Q and Wessels B W 1993 J. Vac. Sci. Technol. B 11 1652 [13] Su Y K, Juang F S and Su C H 1992 J. Appl. Phys. 71 1368 [14] Kuramochi E K and Takanashi Y 1995 J. Appl. Phys. 77 5706 [15] Baranov A N, Cuminal Y, Boissier G, Nicolas J C, Lazzari J L, Alibert C and Joulli´e A 1996 Semicond. Sci. Technol. 11 1185 [16] Ferguson I T, de Oliveira A G and Joyce B A 1992 J. Crystal Growth. 121 267 [17] Michel E, Singh G, Silvken S, Besikci C, Bove P, Ferguson I and Razeghi M 1994 Appl. Phys. Lett. 65 3338 [18] Bertru N, Nouaoura M, Bonnet J, Lassabatere L, Bedel E and Munoz-Yague A 1992 Appl. Surf. Sci. 68 357 [19] Kuramochi E, Kondo N, Takanashi Y and Fujimoto M 1993 Appl. Phys. Lett. 63 2664 [20] Rouillard Y, Lambert B, Toudic Y, Baudet M and Gauneau M 1995 J. Crystal Growth. 156 30 [21] Subanna S, Tuttle G and Kroemer H 1988 J. Electron. Mater. 17 297 [22] Longenbach K F and Wang W I 1991 Appl. Phys. Lett. 59 2247 [23] Lee C D, Nicholas D J, Singer K E and Hamilton B 1986 J. Appl. Phys. 59 2895 [24] Bata S, Horita H and Kinbara A 1978 J. Appl. Phys. 49 3632 [25] Adachi S 1987 J. Appl. Phys. 61 4869 [26] Shin J, Hsu Y, Hsu T C, Stringfellow G B and Gedridge R W J 1995 Electron. Mater. 24 1563 [27] Bertru N, Brandt O, Klann R, Mazuelas A, Ulrici W and Ploog K H 1997 Phys. Rev. B 69 4503