Near-equilibrium LPE growth of GaAs-Ga1−xAlxAs double heterostructures

Journal of Crystal Growth 27 (1974) 86-96 9 North-llolland Publishhtg Co.

NEAR-EQUILIBRIUM LPE GROWTH OF GaAs-Ga~_~AI~As DOUBLE HETEROSTRUCTURES L. R. DAWSON

Bell Laboratories, ;~lurray Hill, New Jersey 07974, U.S.A. Received 13 March 1974; revised manuscript received 22 May 1974 A system for the reproducible liquid phase epitaxial (LPE) growth o f GaAs-Ga~_xAI~,As double heterostructures has been developed. The contamination o f Al-bearing growth solutions with oxygen, and the consequent formation of AIzO3, has been eliminated by baking the Ga under Hz prior to the addition o f AI. This procedure permits uniform wettingof the substrate by the melt and complete removal of the solution to terminate growth. By using a source wafer o f GaAs to seed the solutions prior to arrival o f the substrate, a condition of local thermal equilibrium is maintained at the growth interface at all times. Using these techniques we have reproducibly grown layers 0.15 lam thick. We have applied an isothermal diffusion model to simulate the growth of GaAs from a Ga-As solution. The results of this simulation verify the importance o f maintaining local equilibrium near the growth interface. Using this simulation we have measured the diffusivity of As in Ga, D, and find D = 8 4- l x l0 -5 cm2/sec at 800 ~

1. Introduction

Recent developments in the liquid phase epitaxial (LPE) growth of GaAs-Ga~_~AI~As double heterostructures have allowed significant improvements in the performance of injection lasers. Threshold current densities of --, 1000 A/cm 2 or lower ~-a) and continuous operation at room temperature and above 3"4) have been reported. The essential elements of a typical double heterostructure (DH) used for laser fabricaiion are included in fig. 1. In this configuration, under forward bias, electrons injected from layer [ into layer II (the active region of the laser) are confined there by the energy barrier of the wider bandgap material of layer III. Such carrier confinement enhances population inversion in the active region at a given current densityS). In addition, the lower refractive index of the two Ga~_~AI~As layers surrounding the active region optically confines most of the emitted radiation within the active region6). Both types of confinement, electrical and optical, increase the probability of stimulated emission at a given current density. To date, these structures continue to be prepared exclusively by growth from the liquid rather than from the vapor phase, principally because the controlled vapor transport of highly reactive AI presents severe problems. The presence of AI causes a number of problems in LPE growth as well. Residual oxygen in

the melt or gas stream reacts with AI to form solid AI203, which is not reduced by H z at temperatures normally used for LPE. This A1203 exists as a skin on the melt surface which can seriously interfere with both the uniform wetting of the substrate by the melt and the complete removal of the melt to terminate growth. Such problems probably have been responsible for many of the difficulties experienced in DH growth in the past (refs. 1,7, 8). In the system reported here the dominant 9source of oxygen has been found to be that contained in the Ga, the majority constituent ofall growth solutions. By baking the Ga under H2 prior to the addition of AI this source of oxygen and its consequent A!203 contamination have been virtually eliminated. This procedure results in uniform substrate wetting and removal of the

P GaAs (Ge)

I~Z-,~ l ~ m

P Gaj_x AI x As (Ge)

Ilia- I/~m

P GoAs (Ge)

11 0 1 5 - IO/Jm ACTIVE

n Got_x A i x As (Te or Sn)

n GoAs SUBSTRATE

Fig. 1.

86

Double heterostructure for laser fabrication.

NEAR-EQUILIBRIUM L P E GROWTtt OF GaAs-Ga~_~,AI~,As DOUBLE tIETEROSTRLICTURES

melt from > 95 ~o of the surface after termination of growth. in DH growth, problems are also encountered in the control of compositional and dimensional uniformity. Again, the addition of Al to the Ga-As system poses a problem, in this case, due to the very large segregation of A1 in favor of the solid. In the system reported here the distribution coefficient orAl, knh is about 76. Such a large kn~ can cause the amount of AI incorporated into the growth layer to decrease markedly as growth proceed s due to local depletion of Al from the solution adjacent to the growing interface. This problem is accentuated by high growth rates and/or macroscopic depletion of A! from the bulk of the melt during the growth of thick layers. The severe dimensional tolerances indicated in fig. 1 stem in part from the desire to obtain low values of threshold current density, J,h. Since Hayashi et al. 6) report a nearly linear dependence of Jth on active layer width, d, with Jth (A/cm2) ~ 5 • 10 3 d (lam), values of d as small as 0.15 lain are of interest. The interface bounding this region must be free of protrusions or other irregularities which can scatter light out of the cavity or locally decrease the width d. Additional severe dimensional tolerances exist for layers III and IV due to the requirement that the heat generated within the device be wholly dissipated through a heat sink contacting the last grown layer. This requirement is made more severe by the lower thermal conductivity of Gcq_xAl~,As relative to that of GaAs for values of x in the range of interestg). LPE systems previ9usly used for DH growth have permitted significant departure from equilibrium within the growth solutionst'7). It is shown below that under practical growth conditions from relatively thick melts a considerable degree of supersaturation within the melt is unavoidable. During the sequential growth of numerous layers the substrate may contact solutions which are quite supercooled, causing very large initial growth rates and, consequently, difficulty in controlling very small layer thicknesses. In this paper it is shown that by properly seeding the growth solution with a wafer of GaAs source material prior to contact with the GaAs substrate, the growth solution can be held in local thermodynamic equilibrium in the vicinity of the eventual growth interface. Combining this 'nearequilibrium' technique with low cooling rates (0.1 ~ we have very reproducibly obtained active

87

layer thicknesses of 0.15 + 0.02 lain with interface irregularities less than 0.02 lam. We attribute much of this capability to the use of the source wafer and its effect on the supercooling of the melt prior to the arrival of the substrate. To understand the nature and magnitude of such effects the LPE growth of GaAs from Ga solution has been mathematically simulated under the assumption that As is transported to the growth interface by diffusion alone. The details of this simulation are described by Rode t~ and some of the results are applied to the system reported here. Before describing the growth system we discuss some of the more general aspects of diffusion-controlled LPE, principally to illustrate more fully the need for the source wafer. 2. Diffusion controlled equilibrium conditions Consider the simple growth system shown schematically in fig. 2a. Assume that the system exists at a uniform temperature TA and that the Ga solution has been saturated with As, that is, the melt composition lies on theGa-rich side of the Ga-As iiquidus line at point A of fig. 2b. Moving the GaAs substrate beneath the melt establishes a stable interface between the saturated solution and the solid. Cooling the system uniformly to TB produces supersaturation, creating

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As SUBSTRATE (a) 1238 ~

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~ TB ~ 800 Q_

400

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I I I ] 29.8~ 0 Xc XA Go

50

100

ATOMIC PERCENT ARSENIC

(b) Fig. 2. (a) Single layer growth system. (b) Schematic Ga-As phase diagram; the liquidus has been greatly distorted for illustrative purposes.

88

L.R. DAWSON

uniformly within the melt "excess As", which we define as that concentration of As present in excess of the liquidus composition. In the vicinity of the substrate this departure from equilibrium is relieved by the epitaxial precipitation of GaAs. Assuming that equilibrium is maintained at the interface is equivalent to assuming that each excess As atom which reaches the interface captures a Ga atom and precipitates epitaxially. For a temperature of 800 ~ the distribution coefficient of As is about 20, hence As is rapidly depleted from the solution adjacent to the growing interface. For growth to be sustained, and for equilibrium to be approached by the entire melt, As must be transported to the interface from the bulk of the melt by diffusion and/or convection. We Wish to examine here the case in which As is transported predominantly by diffusion. For finite values of the As diffusivity, D, excess As cannot diffuse to the interface fast enough to maintain equilibrium everywhere within the melt. The distance into the melt over which precipitation at the substrate is moderately effective in maintaining equilibrium can be approximated as the diffusion length of As in liquid Ga, 1:

/ = (ht)~. Taking D = 8 x 10- s cm2/sec for T = 800 ~ a value which we have now measured (see below),.and a growth time t = 20 min, 1 = 3.1 mm. Thus, for a melt 7 mm deep (a 5 g Ga column above a 1.25 cm 2 area substrate), only about 45 7o of the melt participates significantly in the growth process. Since I oc t 89this region of influence increases slowly with growth time. To provide quantitative information on the flux of As atoms within the melt we have simulated LPE growth by a one-dimensional isothermal diffusion model, assuming that the liquid-solid interface remains in equilibrium and no homogeneous nucleation occurs within the melt. The details of this calculation are reported elsewhere t o). From this simulation we obtain the thickness of the grown layer, d, as a function of a number of parameters including the melt thickness, IV, the initial temperature, T o, the cooling rate, 2P, the As diffusivity, D, and the growth time, t. In addition we obtain both the concentration of excess As as a function of position within the melt, and the effective melt fraction, which we define as that fraction of the total excess As created by cooling which has been incorporat-

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50

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I

i. 7

6 5 4 3 2 I DISTANCE FROM GROWTH [NTERFACE(rnrn) ~

O SUBST~ATE

Fig. 3. T i m e development o f the excess As within the melt for growth initiated from equilibrium.

ed into the grown layer. Consider first the simulation of a growth cycle similar to that used in the system which is reported below. Beginning at time zero, with the substrate in equilibrium with an entire 7 mm deep solution, the system is cooled from 800 ~ at a constant rate of 0.1 ~ In fig. 3 a family ofcurves shows the time development of the local excess As concentration within the melt. For convenience the excess As concentration has been represented by an equivalent amount of supercooling. After an elapsed time of 5 min, or 0.5 ~ cooling, nearly the entire melt is supercooled by 0.5 ~ which is consistent with a diffusion length of 1.5 mm for this growth time. At this time the melt fraction is only 0.17. That is, only 17 7o of the excess As existing at this time has diffused to the growth interface. As cooling proceeds the excess As gradient increases, with a consequent increase in As flux toward the growth interface, causing an increasing growth rate, as shown in fig. 4. In the layer thickness curve of fig. 4 we include the dashed line for growth from a melt 3.5 mm deep. This clearly indicates that only near the end of the cycle does the top half of the melt begin to contribute measurably to the growth. This growth system can be extended to multilayer growth by adding an additional melt adjacent to the existing one. If we begin with both melts saturated and proceed as above, the second melt will become uniformly supersaturated by an amount equivalent to the cooling used to grow the first layer. Subsequent movement of the substrate will bring it into contact with this

NEAR-EQUILIBRIUM LPE GROWTII OF GaAs-GaI_:,Ai~As DOUBLE i I E T E R O S T R U C T U R E S i

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Fig. 4. Growth rate and layer thickness during the cooling cycle. supersaturated melt and the growth interface will not initially be at equilibrium. This leads to an extremely fast initial growth rate. Fig. 5 shows the layer thickness versus growth time for various degrees of initial uniform supersaturation. For AT(first layer) = 4 ~ as is typical in the procedure reported below, the first minute of growth produces a layer more than 1 lain thick. For the growth of layers ~ 0.2 I~m thick only 2 sec is required. Even for ATas small as 2 ~ only 5 sec 5,

4

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is required. It is dear that allowing such departure from equilibrium at the liquid-solid interface makes it very difficult to control the growth of layers as thin as 0.2 pro. In the system reported here we use a GaAs source wafer, shown in fig. 6, to seed the second solution during the growth of the first layer. Growth on this source wafer maintains local equilibrium at the bottom of the second melt so that when moved, the substrate does not contact a severely supersaturated region of the solution. We can also use this simulation to examine the effects of growth from undersaturated solutions. Consider the substrate in the position shown in fig. 2, but with the solution approximately 1% deficient in As relative to the 800 ~ liquidus composition, which coresponds to - l ~ supercooling. When the substrate is moved beneath the melt and a cooling rate of 0.I ~ is simultaneously initiated, the substrate becomes a source of As and is initially dissolved to bring the melt to the liquidus composition adjacent to the growth interface. Arsenic (GaAs) continues to dissolve and diffuse into

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Growth rate during the cooling cycle for various of initial supersaturation.

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Fig. 7. Time development of the excess As within the melt for growth initiated from an undersaturated solution.

90

L.R. DAWSON I

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undersaturated solution. the melt until the system has cooled enough to establish a positive excess As concentration gradient at the interface. In fig. 7 we see that not until 5 min has elapsed does this slope become positive. Fig. 8 shows that dissolution of the substrate occurs for about 5 min, during which 0.62 pm of the substrate is consumed. Growth actually begins at about t = 5 min, but the interface is not restored to its initial position until t . = 15 rain. If the melt were undersaturated by an amount equivalent to - 3 ~ supercooling (N 3 % As deficient) the initial meltback time would be 15 min, during which 3.23 pm would be consumed. Three important results are evident here. First, if a relatively thin layer had been previously grown on the substrate, a significant part of it would be lost. Second, the layer grown during this cycle would be considerably thinner than anticipated for an assumption of initial equilibrium and would, in fact, be strongly dependent upon the exact degree of undersaturation. For the 1% As deficient example, in 15 min the growth would be 0.62 I-tm, as opposed to 1.59 Ilm had equilibrium existed at the beginning of the cycle. Ofcourse, the attempted growth of a layer <0.62 lam would result in no layer at all. Third, not all of the dissolved material will have diffused away from the interface before growth finally begins, and some undesirable constituents may be incorporated into the grown layer. These significant effects, Which arise from only a modest degree of undersaturation, are minimized

Time

needed for saturation.

by the presence of the source wafer before the arrival of the substrate. Since precise saturation is desirable we have simulated the process o f heating a melt rapidly from room temperature to 800 ~ while in contact with GaAs, and examined the effective melt fraction while holding the temperature constant. The results shown in fig. 9 indicate that 2 hr must elapse before the amount o f As in solution reaches 99 % of that needed for saturation. In practice, the saturation time may not be this large, since in the initial stages of rapid heating, when dissolution of GaAs is most rapid, there is undoubtedly some corivection due to initial thermal nonuniformities. After thermal stabilization the solution will be richer in As near the dissolving source GaAs, and may be slightly more dense there, since 3lAs > 3/Ga. If this source GaAs is floating on the top surface of the melt, as is the case in many growth systems, the slight density gradient may cause some convection. On the time scale considered here, even a small amount of convection could be significant. By using the source wafer to aid in saturation from the bottom of the melt in a manner similar to that of Lockwood and Ettenberg11), and Blum and Shih12), the above time requirements are somewhat relaxed.

3. Experimental The essential elements of the growth system are shown in fig. I0. The boat is positioned within the quartz envelope by means of a quartz rod. The slider is adjusted by a push rod. In the cold section of the envelope a loading tube is placed so that the boat can be withdrawn from the furnace and AI, GaAs source

N E A R - E Q U I L I B R I U M L P E G R O W T H OF

GaAs-Ga~_xAl~As DOUBLE I I E T E R O S T R U C T U R E S

H2 EXHAUST .~ SLIDER PUSH ROD

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POSITION

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THERMOCOUPLES

Fig. 10. Growth apparatus.

material, and dopants can be added to each of the growth solutions without significant back diffusion of air. Pd-diffused H 2 enters at one end of the tube and exits at the other end, near the necessary sliding seals. The exhaust end, from which the system is loaded, is contained within a dry N2 ambient so that loading can be done safely without interruption of the H2 flow. This constant H2 environment within the envelope minimizes problems associated with the adsorption of other gases, particularly oxygen, on the system walls. In the 3-zone furnace the end winding control thermocouples are slaved to the center thermocouplq so that no thermal gradients will be introduced during the cooling cycle. Three probe thermocouples are located in fixed positions beneath the boat. Their spacing permits measurement of the temperature gradient along the length of the boat as well as the boat temperature during the entire cooling cycle. Using this arrangement it is possible to maintain temperature inhomogeneities _< 0.5 ~ (at 800 ~ along the length of the boat (11 cm) during all phases of the growth of a 4-layer double heterostructure. Under constant temperature conditions the center zone maintains a long term (,,- 3 hr) thermal stability of _+0.05 ~ Cooling is effected by introducing an electronically generated ramp voltage in series with the center winding control thermocouple. The graphite boat assembly, fig. 6, consists of a body having fourcompartments (crucibles) ofequalsize and spacing and a slider having two similarly spaced recesses for the substrate and source wafers. All parts are machined from type DFP-! Poco graphite, which has isotropic thermal conductivity and a nominal im-

purity level of 200 ppm before machining. Before use, all mating surfaces are lapped with 1 lam alumina abrasive to remove machine marks and to facilitate the removal of loosely adhered graphite dust. The procedure for the growth of a four-layer double heterostructure similar to that shown in fig. 1 begins by placing an undoped, polycrystalline GaAs source wafer and a <100> n-type GaAs substrate into the slider. The substrate has previously been chemically polished in a Br-methanol mixture and just before use is cleaned in organic solvents (chloroform and methanol) and dilute aqueous HCI, finishing with direct decanting into methanol and drainage onto filter paper. The substrate is loaded on top of a graphite spacer thick enough to bring the substrate surface within ,-, 75 pm of the slider surface (and hence the wiping edge). This spacing has not been optimized, but 75 lam gives very consistent results. Four-layer structures have been grown with complete melt removal using a spacing of ,-, 225 lam. The slider is moved to the position shown in fig. 6, where the substrate is tightly covered by the body of the boat with little space available to receive volatile decomposition products. The crucibles are loaded with appropriate amounts of 6N (99.9999 %) pure Ga and 6N pure Sn. IfTe is to be used as the n-type dopant it is added at this point and covered with Ga to prevent vapor loss before dissolution in the liquid. The boat is then loaded into the envelope and positioned at the center of the furnace ( ~ 800 ~ where it remains for one hour. The purpose of this very hllportallt step is the removal of oxygen contained in the Ga (and Sn) by hydrogen reduction and/or volatilization of Ga20. The

92

L.R. DAWSON

boat is then wiihdrawn from the furnace to a location beneath the loading tube, through which At, source GaAs, and dopants are added to the melts. During this procedure H 2 exhausts through the loading tube, ensuring that no significant back-diffusion of air occurs. The boat is returned to the furnace where a minimum of 3 hr is allowed for equilibration. The slider is moved to contact the source wafer with melt I, where it remains for ~ 30 rain to ensure that equilibrium has been attained. The slider is then moved one ceil length, a, and held for ,-~ 10 min to allow any contaminants which might remain on the substrate surface to be dissolved into the melt. The system is cooled at a uniform rate of 0.1 ~ which is maintained for the duration of the run. Growth occurs simultaneously on the source and substrates, as described above, producing excess As gradients in the growth solutions, but maintaining local equilibrium at the liquid-solid interfaces. When the first layer growth is terminated by moving the slider another cell length, the substrate contacts a region of melt II that is not appreciably supersaturated relative to the bulk of the solution, as shown by results presented below. The process is repeated for the growth of 4 layers, each time with the substrate contacting a liquid of low local supersaturation. During part of the cooling cycle solutions II1 and IV have been cooled without the presence of the source wafer. The only means of maintaining equilibrium is growth on the rather" limited area of any remaining GaAs source material floating on the top of the melt and/or homogeneous nucleation within the melt. The occurrence of homogeneous nucleation is doubtful in view of the fact that immediately after withdrawal from the furnace the surfaces of all melts are liquid and very shiny, showing no evidence of surface nucleation. This observation is consistent with reports 13'14) that similar III-V solutions will supercool considerably more than the few ~ required here for the growth of all layers. 4. Results

In virtually all growth runs using this procedure the melt has been cleanly removed from at least 95 ~o of the surface of the final layers. No evidence of Al2Oa formation is found in any of the growth solutions. Fig. 1la shows the surface of a typical 4-layer DH as its appears when removed from the growth apparatus. The growth area is 1.25 cm • 1.0 cm. Such surfaces are of mirror

Fig. I I. Surfacequality of a typical double heterostructure as removed from the growth apparatus. finish to the unaided eye and show the structure of fig. l lb when inspected with interference contrast microscopy. The uniformity of these layers is shown in fig. t2, a photomicrograph of an angle-lapped DH. The difference in reflectivity between GaAs and Gao.vsAIo.2sAs gives excellent delineation of the interface without staining. The active region of this DH is 0.l 6 lain thick. The irregularities in the interfaces bounding this region are estimated to be less than 200 A. The reproducibility from run-to-run was investigated by growing a series of 5 consecutive DH's under nearly identical conditions. Table 1 shows the thickness and growth time for each layer, and table 2 shows the liquid and solid compositions, For these structures the substrate was in contact with solution II for only I rain, resulting in

NEAR-EQUILIBRIUM L P E GROWTH OF G a A s - G a t _ ~ A l x A s DOUBLE tlETEROSTRUCTURES

~

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1.83/.Lm

f

P-- GOo.TsAs 0.55 ~ m

93

SURFACE p -- Go As

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ACTIVE LAYER p -- GoAS O.16~m N -GOo.66AQ.o.34As 2.26~m

into c o n t a c t with melt I at equilibrium, as described above. A c t u a l l y , a b i n a r y solid ( G a A s ) c a n n o t be in equilibrium with a q u a t e r n a r y liquid ( G a - A I - S n - A s ) , so p r e s u m a b l y a very thin layer o f the necessary solid (dictated by the d i s t r i b u t i o n coefficients o f the v a r i o u s 9 .

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9

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TABLE 1 Layer thickness Sample # 9 -

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I d (lam)/t(min)

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Fig. 12. Photomicrograph of an angle-lapped 4-1ayer double heterostructhre. active region thicknesses in the range 0.15 _+ 0.02 ~tm. T h e u n i f o r m i t y o f Ai c o n c e n t r a t i o n across the thickness o f the G a l _ x A l ~ A s layers was d e d u c e d f r o m reflectivity m e a s u r e m e n t s ~5) on a n g l e - l a p p e d surfaces. T h e v a r i a t i o n o f x in the direction o f g r o w t h in all samples e x a m i n e d was below the detectability limit o f the m e t h o d (+__ 0.02) for layers as thick as 6 lain. N o AI is detected in layers II a n d IV, indicating that no significant a m o u n t o f either A l - b e a r i n g s o l u t i o n is transferred to an a d j a c e n t solution d u r i n g the m o v e m e n t o f the slider. In fig. 12 it is evident that the interface between the substrate a n d layer 1 is very rough. This roughness is virtually always present when we bring the s u b s t r a t e

B 091 B 092 B 093 B 094 B 095

2.20/44 2.35/41 2.33/45 2.11/42 1.96/42

Layer # ll(Active) Ill d/t d/t 0.14/1 0.13/1 0.15/1 0.17/1 0.14/I

0.62/I0 0.78/15 0.97/15 0.87/15 0.52/15

1V d/t 1.55/15 1.63/20 1.59/20 2.21/20 1.80/22

d = layer thickness; accuracy is 4-10%. t = time that substrate remains in contact with melt. TABLE 2 Ga~_~AIxAs composition and electrical data Layer

Ga (at%) Sn (at%) Ge (at%) AI (at%) As (at%)

1

I1

83.6 14.7 0 0.20 i.55

97.6 0 0.13 0 2.23

III

Liquid 97.7 0 0.45 0.15 1.67

IV 97.3 0 0.45 0 2.23

Solid x

. (cm- 3) p (cm -3)

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0.25

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94

L . R . DAWSON

species) crystallizes on the substrate. We find that if we begin the cooling cycle one step earlier, when the source wafer is in contact with melt I, growth is initiated immediately upon arrival of the substrate. This slight modification .consistently produces very smooth initial interfaces. It Should be emphasized, however, that this procedure may trap any contaminants remaining on the substrate surface. Since this interface is only ~ 2 lam from the active region of the laser, minute amounts of undesirable impurities may be important in the degradation of D H devices. Background contamination by electrically active impurities was determined by 300 ~ and 77 ~ Hall measurements made on undoped single layers grown on Cr-doped substrates. We found ND--NA = 4 • 1015 cm-aandNo+Nh ~ 1016cm -3. The thickness of the active layer, d~t, for this series of runs is in very good agreement with the growth simulation described above, assuming that no stirring of the melt occurs when the substrate is moved from melt I to melt II. In this case growth from melt II has occurred on the source wafer for ~ 42 min during the growth of layer I (from melt I) on the substrate. Under the assumption of no stirring, the growth of layer II is just that which would occur during the next minute of growth on the source wafer. For the conditions given here, that additional minute of growth provides a simulated thickness of 0.25 lam, in substantial agreement with the experimental value of 0.15 _ 0.02 ~tm. Any stirring will increase the thickness of the grown layer. If one assumes that the movement of the slider causes the melt to rotate, a region of nearly uniformly supersaturated liquid will contact the substrate. Taking the degree of supersaturation to be 2 ~ a layer 0.84 lam thick would be grown during the first minute. These results strongly indicate that little or no stirring of the melt occurs upon movement of the slider. This is in agreement with the approximation of Rode t~ that a layer of solution only ~ 25 lam thick is disturbed by movement under these conditions. Without the source wafer the growth of all layers is initiated from nonequilibrium interface conditions, leading to much higher growth rates, as discussed above. Fig. 13a shows the surface of a 4-layer D H grown by a procedure similar to that above, but without a source wafer. This rough surface shows obvious signs of the 4-fold symmetry of the (100) orientation, and this

j ..........................................

Fig. 13. (a) Surface of a 4-layer double heterostructure grown without the source wafer. (b) Angle-lapped section of the same wafer. roughness is reflected in the interfaces shown in the angle-lap of fig. 13b. It is not known whether this surface structure results from constitutional supercooling16), irregular surface nucleation, or other growth rate dependent effects. We can only conclude that the omission of the source wafer, and the ensuing lack of local equilibrium causes serious morphology effects. This system, together with the isothermal diffusion model simulation ~o) described above, has been used to measure the diffusion coefficient of As in Ga at elevated temperatures. Single layers of GaAs have been grown on (100) substrates after carefully establishing equilibrium conditions by prolonged exposure of the growth solution to the source wafer. The solid lines in fig. 14 are the simulated layer thickness vs. growth time curves for various values of the arsenic diffusivity at 800 ~

GaAs-Ga~_xAl~,As DOUBLE I t E T E R O S T R U C T U R E S

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The error bars on the data points reflect the range of layer thickness measured at four points on the surface of the wafer. The fact that these data closely follow the shape of the simulated curves and do not show additional curvature upwards is a strong indication that convection does not contribute significantly to the layer thickness. From these data, D is found to be 8 +_ i • 10- s cm2/sec at 800 ~ Fabry-Perot diodes fabricated from DH's prepared in this manner have shown current thresholds, Jth as low as --, 1200 A/cm 2 (d = 0.15 pm), and external differential quantum efficiency, II, as high as 43 %. A representative intensity versus current density curve is shown in fig. 15. 5. Conclusions In this paper we have described an LPE growth system which can be used to reproducibly grow layers as thin as 0.15 ~tm. We have shown that the problems caused by AI203 formation and its detrimental effects on solution-substrate contact, can be averted by removing the oxygen from the solution prior to the addition of AI. We have also illustrated the desirability of maintaining local equilibrium near the growth interface and have shown that seeding by the source wafer provides this function. The degree of control in this system and the apparent lack of significant convection w.ithin the melt have enabled us to accurately measure the diffusivity of As in Ga at 800 ~

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Acknowledgements The author wishes to thank Mrs. D. R. Ketchow for her assistance in the growth of many of the heterostructures, and D. L. Rode for many helpful discussions and the use of his diffusion calculations. References 1) Zh. I. Alferov, V. M. Andreyev, V. I. Korol'kov, E. L. Portnoi and D. N. Treryakov, Kristall und Technik 4 (1969) 495. 2) B. I. Miller, E. Pinkas, I. Hayashi, P. W. Foy and R. Capik, Appl. Phys. Letters 19 (1971) 340. 3) I. Hayashi, M. B. Panish, P. W. Foy and S. Sumski, Appl. Phys. Letters 17 (1970) 109. 4) J. E. Ripper, J. C. Dyment, L. A. D'Asaro and T. L. Paoli, Appl. Phys. Letters 18 (1971) 155. 5) I. Hayashi, M. B. Panish and P. W. Foy, IEEE J. Quantum Electron. 5 (1969) 211. 6) I. Hayashi, M. B. Panish and F. K. Reinhart, J. Appl. Phys. 42 (1971) 1929. 7) M. B. Panish, S. Sumski and I. Hayashi, Met. Trans 2 (1971) 795.

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8) M. B. Panish, I. Hayashi and S. Stimski, IEEE J. Quantum Electron. 5 (1969) 210. 9) M. A. Afromowitz, J. AppL Phys. 44 (1973) 1292. 10) D. L. Rode, J. Crystal Growth 20 (1973) 13. i l ) H. F. Lockwood and M. Ettenberg, J. Crystal Growth 15 (1972) 81.

12) J. M. Blum and K. K. Shih, J. Appl. Phys. 43 (1972) 1394. 13) M. B. Panish, J. Chem. Thermodynamics 2 (1970) 319. I4) M. B. Panish and J. R. Arthur, J. Chem. Thermodynamics 2 (1970) 299. 15) M.A. Afromowitz and R. L. Brown, private communication. 16) H.T. Mindcn, J. Crystal Growth 6 (1970) 228.