High-Performance Near-IR Photodiodes: A Novel ... - IEEE Xplore

IEEE JOURNAL OF QUANTUM ELECTRONICS, VOL. 47, NO. 2, FEBRUARY 2011

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High-Performance Near-IR Photodiodes: A Novel Chemistry-Based Approach to Ge and Ge–Sn Devices Integrated on Silicon Radek Roucka, Jay Mathews, Change Weng, Richard Beeler, John Tolle, José Menéndez, and John Kouvetakis

Abstract— Ge/Si heterostructure diodes based on n++ Si(100)/ i-Ge/p-Ge and p++ Si(100)/i-Ge/n-Ge stacks and intrinsic region thickness of ∼350 and ∼900 nm, respectively, were fabricated using a specially developed synthesis protocol that allows unprecedented control of film microstructure, morphology, and purity at complementary metal–oxide–semiconductor compatible conditions. From a growth and doping perspective, a main advantage of our inherently low-temperature (390 ºC) softchemistry approach is that all high-energy processing steps are circumvented. Current–voltage measurements of circular mesas (60–250 µm in diameter) show dark current densities as low as 6 × 10−3 A/cm2 at −1 V bias, which is clearly improved over devices fabricated under low thermal budgets using traditional Ge deposition techniques. Spectral photocurrent measurements indicate external quantum efficiencies between 30 and 60% of the maximum theoretical value at zero bias, and approaching full collection efficiency at high reverse biases. The above Ge devices are compared to analogous low-temperature-grown (350 ºC) Ge0.98 Sn0.02 diodes. The latter display much higher dark currents but also higher collection efficiencies close to 70% at zero bias. Moreover, the quantum efficiency of these Ge0.98 Sn0.02 diodes remains strong at wavelengths longer than 1550 nm out to 1750 nm due to the reduced band gap of the alloy relative to Ge. Index Terms— Germanium–tin alloys, infrared detectors, integrated optoelectronics, p-i-n, photodiodes, photovoltaic cell materials, semiconductor epitaxial materials, ultrahigh vacuum chemical vapor deposition.

I. I NTRODUCTION

S

EMICONDUCTOR near-IR photodiodes typically used in state-of-the-art fiber network technologies currently

Manuscript received March 8, 2010; revised August 11, 2010; accepted September 5, 2010. Date of current version January 19, 2011. This work was supported in part by the Air Force Office of Scientific Research, Multidisciplinary University Research Initiative, under Grant FA9550-06-01-0442, by the Department of Energy, under Grant DE-FG36-08GO18003, and by the Interconnect Focus Center-Semiconductor Research Corporation/Defense Advanced Research Projects Agency Focus Center, under Task 674.015. R. Roucka, R. Beeler, J. Tolle, and J. Kouvetakis are with the Department of Chemistry and Biochemistry, Arizona State University, Tempe, AZ 85287-1604 USA (e-mail: [email protected]; [email protected]; [email protected]; [email protected]). C. Weng was with the Department of Chemistry and Biochemistry, Arizona State University, Tempe, AZ 85287-1604 USA. She is now with Intel Corporation, Santa Clara, CA 95054-1549 USA (e-mail: [email protected]). J. Mathews and J. Menéndez are with the Department of Physics, Arizona State University, Tempe, AZ 85287-1504 USA (e-mail: [email protected]; [email protected]). Color versions of one or more of the figures in this paper are available online at http://ieeexplore.ieee.org. Digital Object Identifier 10.1109/JQE.2010.2077273

involve expensive III–V materials, which are difficult to integrate with conventional electronics since this requires the application of complex heterogeneous processes. An alternative approach is to use silicon photonic technologies based on ubiquitous group IV materials such as elemental Ge and Ge1−x Six alloys monolithically integrated with conventional electronic functions. However, a major limitation of Ge1−x Six alloys is that their bandgaps do not cover the entire telecom range between 1260 and 1670 nm. Even elemental Ge barely reaches the 1550-nm wavelength required for the so-called C-band window. Another disadvantage is that their inherent lattice mismatch with the underlying Si platform produces structural and morphological imperfections, which hinder the IR detector performance. Although defect-free Ge1−x Six layers can be grown pseudomorphically on Si wafers, their critical thickness must be reduced to unacceptably low levels in order to bring their optical absorption edge closer to the needed IR range by increasing the Ge fraction in the alloy. Given the limitations of Ge1−x Six alloys, recent efforts have focused on developing non-pseudomorphic Ge detectors on Si by systematically improving the microstructural quality of the active layers. In this regard, substantial progress has been made over the past 10 years to minimize the dislocation densities from the typical 108 −109 /cm2 range [1] observed in early studies down to recently reported values as low as 105 −106 /cm2 . This breakthrough is associated with the development of creative synthesis and processing methods, including new templating concepts such as compositionally graded Ge1−x Six buffers [2] and growth of low-temperature Ge initiation layers followed by high-temperature bulk deposition [3]. The resulting materials have been used to fabricate several proof-of-concept photodetectors, monolithically integrated with Si electronics. However, the required thermal budget to fabricate such devices, which in some cases involve high-temperature annealing up to 900 ºC [4], is often incompatible with complementary metal–oxide–semiconductor (CMOS) process flows. More recently, efforts have been made to lower the growth temperature and to eliminate annealing steps. p-Si/i-Ge/n-Ge heterojunction p-i-n devices with intrinsic layer thicknesses of 800 nm grown at temperatures not higher than 600 ºC and without additional annealing show dark-current densities of 0.1 A/cm2 at a reverse bias of 1 V and an unbiased responsivity of 0.4 A/W at 1300 nm and 0.2 A/W at 1550 nm [2]. The inferior responsivity of Ge around the

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L telecom window (1560–1620 nm) has also been addressed by inducing tensile strain on the Ge layers, which lowers the direct bandgap [4]. The amount of tensile strain, however, is roughly proportional to the growth or annealing temperature, and therefore high levels of tensile strain are incompatible with low thermal budgets. So far, the lowest reported growth temperature to achieve high tensile strain was 600 ºC, but it required the growth of a double Si/SiGe buffer layer prior to Ge deposition [3]. Recently, we have developed an entirely new lowtemperature approach to the growth of Ge on Si [5], which circumvents the well-known Stranski–Krastanov growth mechanism. This method uses a metalorganic additive at ∼390 ºC, and yields films with extremely low dislocation densities (104 cm−2 ), atomically flat surfaces, and thicknesses approaching 5 μm. Since the growth temperature of Ge on Si (390 ºC) under this method is substantially reduced compared to conventional approaches, our Ge layers are singularly attractive for the integration of near-IR detectors with Si under CMOS-compatible conditions. In this paper, we present Si/Ge heterojunction p-i-n diodes fabricated via the new low-temperature route using both n-type and p-type Si substrates. As indicated above, low-temperature growth limits the level of tensile stress that can be applied on the Ge layers, and therefore our metalorganic additive method is not compatible with detection within the telecomm L band (1565–1625 nm) below the bandgap of relaxed Ge. We have recently demonstrated an alternative approach to extend the near-IR sensitivity of Ge-based detectors by incorporating Sn substitutionally into Ge [6], [7]. Only 2 at% Sn is sufficient to produce materials that provide full coverage of all telecom windows. Furthermore, the absorption coefficient of Ge0.98 Sn0.02 at 1550 nm is about 10 times higher than that of Ge, and even at 1300 nm the alloy has a 40% higher absorption coefficient than bulk Ge. This enhancement bodes well for the development of telecom detectors with improved IR performance based on these alloys. From a materials synthesis perspective, the growth of monocrystalline Ge1−y Sn y alloys has been demonstrated directly on Si substrates, and doping protocols of these materials have been optimized to produce high quality n- and p-type layers for subsequent assembly of p-n junctions [8]. Most importantly, all of the intrinsic and doped materials are grown at low temperatures (T < 350 ºC) via reactions of custom-designed Si–Ge–Sn and P(As)–Ge–Si hydrides using conditions that are fully compatible with CMOS back-end processing. Prototype Ge–Sn p-i-n photodiodes showing an extended IR responsivity have been recently demonstrated [9]. In this paper, we expand upon this theme and provide a detailed account of their fabrication and properties. These results are then compared with those of their Ge-based counterparts. We specifically demonstrate that the Ge0.98 Sn0.02 devices suffer from much higher dark currents, but that their optical properties are better, they show significant responsivity at 1700 nm, well past the long-wavelength limit of Ge-based detectors, and at shorter wavelengths they have higher zerobias collection efficiencies than the Ge devices.


hv

Cr/Au

Sio2

Cr/Au

p Ge

i Ge Cr/Au

Cr/Au n+ Silicon

Fig. 1.

Schematic cross-sectional view of the p-i-n photodiode device.

II. R ESULTS AND D ISCUSSIONS A. Growth and Characterization of Ge Diode Structures To minimize the number of Ge layers, we built heterojunction diodes consisting of either an n-type Ge top layer, an intrinsic Ge layer, and a p-type Si substrate (subsequently denoted as n-i-p structure), or a p-type Ge top layer, an intrinsic Ge layer, and an n-type Si substrate (subsequently denoted as p-i-n structure). The schematics of the p-i-n structure is shown in Fig. 1. The Ge stacks are grown on 4-in n-type (As) substrates with a resistivity ρ = 0.003 cm and p-type (B) Si(100) substrates with ρ = 0.02 cm. This corresponds to carrier densities n = 2 × 1019 cm−3 and p = 4 × 1018 cm−3 , respectively. Our newly developed approach to producing the intrinsic and doped Ge device components proceeds via thermal decomposition of digermane (Ge2 H6 ) as the main source of Ge. The digermane is combined with a purposely engineered metalorganic additive – (GeH3 )2 CH2 (digermyl methane) – in a gas-source molecular beam epitaxy (MBE) system at temperatures between 380 and 390 ºC. The growth mechanism is based on the initial adsorption of (GeH3 )2 CH2 to the underlying surface via Ge–Ge bonds. The molecule then serves as a pseudosurfactant, which inhibits the lateral diffusion of the main growth component, i.e., digermane, on the surface [5], thus promoting layer-by-layer growth. It is important to emphasize that the CH2 group from (GeH3 )2 CH2 is released as CH4 . This process is highly efficient, to the extent that high-resolution secondary ion mass spectroscopy (SIMS) depth profiling shows no evidence for the incorporation of any residual carbon traces. In a typical growth experiment, the wafers are cleaned by a modified RCA process with the final step consisting of a 1-min dip in diluted (5%) HF/methanol followed by a methanol rinse. After drying in nitrogen flow, the wafers are introduced into the MBE growth chamber via a load lock, and outgassed at 450 ºC under ultrahigh vacuum. Before deposition, the surface of the wafer is cleaned by a brief flash to 800 ºC for several seconds to remove any possible contaminants. The reflection high-energy electron diffraction

ROUCKA et al.: NOVEL HIGH-PERFORMANCE NEAR-IR PHOTODIODES

1023

At. concentration (/cm3)

1022 Ge Si P

1021 1020 1019 1018 1017

200

400

600 Depth (nm)

800

1000

Fig. 2. SIMS atomic concentration profile showing uniform distribution and abrupt doping profile of P in the top Ge layer of the n-i-p structure. The phosphorus signal in the Si substrate is an artifact caused by an interference of 30 Si-H complex (31 amu).

(RHEED) image of the resultant wafer exhibits a distinctly sharp 2 × 1 pattern, indicating an atomically flat surface free of any adsorbed species. A gaseous precursor mixture of (GeH3 )2 CH2 , Ge2 H6 , and H2 in the ratio 1:20:20 is then introduced into the chamber via a leak valve and dispersed across the wafer surface through a custom-designed shower head to ensure a uniform thickness profile across the 4 in wafer. Growth pressures are maintained at ∼0.1 mTorr at a temperature of 390 ºC to yield layer thicknesses of 250–2000 nm with typical growth rates approaching 6–10 nm/min. The required doped Ge top layers of the device heterostructures are grown under the same conditions by adding appropriate amounts of (GeH3 )3 P (n-type) or B2 H6 (p-type) into the reaction mixture to obtain carrier concentrations in the 1019 cm−3 range. In situ RHEED images of the Ge samples also exhibit a streaky 2 × 1 pattern, indicating an atomically flat surface consistent with layerby-layer growth. Noncontact ultraviolet and IR spectroscopic ellipsometry were used to determine the thicknesses of the intrinsic and doped Ge layers and estimate the doping levels. Diode structures were fabricated starting with an n-i-p sample consisting of a ∼100-nm n = 2 × 1019 cm−3 top layer and a ∼900-nm intrinsic layer, and a p-i-n sample containing a ∼100-nm p = 1 × 1019 cm−3 top layer and a ∼350-nm intrinsic layer. The doping concentrations measured by SIMS and ellipsometry yielded very similar values, indicating full dopant activation. As shown in Fig. 2, the elemental profiles of the n-i-p structure indicate a uniform distribution of the P donors in the top layer, which decays rapidly below the detection level in the vicinity of the interface with the underlying intrinsic Ge. From these data, we conclude that the impurity levels in the intrinsic layers are below the detection limit of ∼1016 cm−3 . Prior Hall measurements of nominally intrinsic films of comparable thicknesses grown on highly resistive Si wafers show that they are n-type with carrier concentrations in the n = 7 × 1015 − 5 × 1016 cm−3 range. After growth, all structures were subjected to three cycles of rapid thermal annealing (RTA) at 680 ºC for 10 s. While

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this does introduce a higher temperature step into the material processing, the short duration of these RTA cycles keeps the thermal budget to a minimum. X-ray diffraction θ/2θ scans revealed a single (004) peak with an ω rocking curve full-width-at-half-maximum (FWHM) ranging from 225 to 450 arcseconds for the thick (900 nm) and thin (350 nm) samples, respectively. To our knowledge, these FWHM values are better than state-of-the-art for Ge-on-Si heteroepitaxy [10]. Furthermore, the annealing step does not cause any discernible intermixing between the film and the substrate at their interface. Similarly, we note that the shape of the profile obtained by the in situ doping protocol remains unchanged after RTA processing of the samples, which proves that the sharp device junction is thermally robust. The lattice parameter obtained from the (004) X-ray diffraction peak agrees with that of bulk Ge, indicating that the grown layers are strain free. The microstructure of the device stacks was also investigated by cross-sectional transmission electron microscopy (XTEM), and representative micrographs of the 350- and 900-nm thick Ge samples are shown in Fig. 3. The results demonstrate flat epitaxial films. In the low-resolution images shown on the left (top/bottom), threading defects are not visible throughout the layer within the entire field of view. The mismatch between Ge and Si is absorbed by a periodic array of Lomer dislocations located at the interface. These are clearly visible in the high-resolution image shown on the top right of Fig. 3, with a spacing corresponding to the ratio of the Si and Ge lattice parameters, as expected for a relaxed film. The smooth morphology is further corroborated by ex situ atomic force microscopy (AFM), which yields rms roughness values of ∼0.3 − 0.4 nm for a 20 μm × 20 μm area, as shown in the bottom right panel of Fig. 3 for the 900-nm thick device stack. B. Ge Device Fabrication The fully characterized Ge samples were processed into photodiodes, which were designed as circular mesas etched down to the surface of the Si substrate. The top and bottom electrical contacts consisted of ring-shaped regions based on 20- and 200-nm thick Cr/Au layers, respectively. The device fabrication steps employed here (photolithography, etching, metal contacts) were developed during our previous studies of Ge–Sn-based photoconductors [11] and p-i-n diodes [9]. Initially, the surface of the Ge film was cleaned by sonicating in methanol, followed by a deposition of a 100-nm silicon oxide layer using plasma-enhanced chemical vapor deposition (PECVD). The latter served as a protection barrier during subsequent processing of the samples. Standard photolithography using AZ3312 resist was employed to pattern the mesas, which were then formed by reactive ion etching using BCl3 plasma at 50-mTorr pressure and 80-sccm flow, resulting in typical etch rates of 50 nm/min. After stripping the photoresist and protective oxide layers, the surface of the Ge film was passivated by dipping the samples into diluted HF/methanol solution for 1 min. This was immediately followed by PECVD deposition of a 420-nm thick silicon oxide coating. The metal contact pads were defined using AZ4330 resist, which was then baked at 80 ºC for 30 min to enhance its negative lip

216


Ge

Ge

Si

200 nm

5 nm

Si

Ge

0.5 µm

5 µm

Si (100)

Fig. 3. XTEM image of the 350- and 900-nm thick Ge films (left top and bottom, respectively) showing a flat surface and no threading defects propagating through the layer. The high-resolution image on the top right shows edge dislocations (arrows) located at the Ge/Si interface. (bottom right) AFM image of an atomically smooth Ge surface exhibiting rms roughness of 0.4 nm.

profile for successful metal liftoff. The silicon oxide layer in the contact areas was etched away by a buffered oxide etchant, and Cr/Au metal contact stacks with thickness 20/200 nm were then deposited immediately thereafter by e-beam evaporation. The formation of electrical pads was finalized by liftoff in acetone, and the resultant devices were finally cleaned in an oxygen plasma. Optical microscopy and scanning electron microscopy (SEM) examinations showed that the mesa definition and lift-off process for the metal contacts were successfully conducted in all samples. This is demonstrated in Fig. 4, showing a plan view image of a 300-μm diameter photodiode structure (including the metal contacts), which is typical of those fabricated in an array of devices with variable sizes ranging from 60 to 300 μm. C. Electrical Measurements The initial electrical characterization of all Ge photodiodes involved measurements of their current—voltage (I–V ) curves. Typical results are shown in Fig. 5 (top) for the p-i-n sample with a 350-nm thick i-Ge layer, and in Fig. 5 (bottom) for the n-i-p sample with a 900-nm i-Ge layer. The plots present a set of curves obtained from diodes with diameters 60, 100, 150, 200, and 250 μm. The dark current densities do not show a strong correlation with the device diameter, from which we conclude that the leakage current is generated at the Si/Ge interface or within the bulk of the material rather than the sidewalls. From the slope of the linear region of the I–V curves (0.0 − 0.2 V), we have extracted an ideality factor of 1.2, which indicates low recombination current, as expected for high-quality active Ge regions. Further evidence for the high quality of our Ge layers is provided by the low dark current values, which compare

bottom contact top contact

SiO2 window

mesa edge 20 μm

Fig. 4. SEM image of a representative Ge photodiode showing a well-defined circular mesa (300 μm in diameter) and the corresponding Cr/Au contact pads.

very favorably with those reported for structures prepared under low thermal budgets. For example, Suh et al. find dark current densities of 2 × 10−2 A/cm2 in p-i-n structures with intrinsic layer thicknesses of 1200 nm grown at 650 ºC without further annealing [12]. Colace et al. report dark current densities of 0.2 A/cm2 in n-i-p structures with 1000-nm intrinsic layers grown at 600 ºC without annealing [2], and Jutzi and coworkers find dark current densities of 0.1 A/cm2 in n-i-p structures with an intrinsic layer thickness of 300 nm and a bottom p-Ge layer grown by MBE at 300 ºC [13]. In fact, our structures with 350-nm intrinsic layers have lower dark currents than diodes with similar intrinsic layers


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103 Dark current density @ −1 V

103 250 μm

101

200 μm 150 μm 100 μm 60 μm

−1

Current Density (A/cm2)

10

102

Colace and Assanto

101 100 10−1

This Ge

10−2 10−3 101

10−3

102

103

107 105 Dislocation density (cm−2)

109

Fig. 6. Dark current versus dislocation density plot showing a best fit through the data of Colace and Assanto and the corresponding extrapolation into the regime of our electrical measurement.

250 μm 200 μm 150 μm 100 μm 60 μm

−10.5

100 −11.0

In (I)

10−2

−11.5 −12.0

10−4 −1.0

−0.5

0.0 Bias (V)

0.5

1.0

Fig. 5. Current density (A/cm2 ) curves for photodiodes with diameters ranging from 60 to 250 μm for samples with intrinsic region thickness of 350 (top) and 900 nm (bottom). The dark current densities span from 6 × 10−3 to 1 × 10−2 A/cm2 , which are on the order of the best results reported to date on similar thickness materials and device geometries, and are independent of the device diameter.

grown at much higher temperatures [14]. We believe that the reason for the observed low dark currents in our devices is the low density of threading dislocations. In Fig. 6, we show the correlation between dark current and dislocation density as noted by Colace and Assanto [1]. The grey box corresponds to our measured dark currents combined with dislocation densities found in similarly grown Ge layers [5]. Our results are in reasonable agreement with an extrapolation of the Colace—Assanto line, particularly if one considers the large errors inherent to an extrapolation of a log–log relationship. Further insight into the origin of the dark current can be obtained by examining its temperature dependence. Arrhenius plots of dark current versus temperature for a set of reversebias values between −0.2 and −1 V are shown in Fig. 7 for the p-i-n sample with the 350-nm intrinsic layer. From the slope of these curves, we obtain thermal activation energies E a , which decrease with applied bias. The low-bias activation energies are noticeably higher than half the Ge bandgap, indicating a significant diffusion contribution to the dark current [15], as expected from the low ideality factor of 1.2. Masini and coworkers [16] have modeled the dark current in n-i-p and p-i-n devices. The p-i-n structure is found to be

−12.5

−1 V −0.8 V −0.6 V −0.4 V −0.2 V

35.0

Ea = 0.32 eV Ea = 0.37 eV Ea = 0.42 eV Ea = 0.46 eV Ea = 0.51 eV

36.0 1/kT eV−1

37.0

Fig. 7. Arrhenius plot of the dark current showing an increase of the activation energy with increasing reverse bias. The activation energies E a > E g /2 indicate that the dark current is dominated by diffusion.

independent of the interface recombination velocity, but for n-i-p structures the dark current depends strongly on the interface recombination velocity for the case of highly doped substrates, such as those used here. Therefore, a direct comparison of dark currents between p-i-n and n-i-p devices is not straightforward. D. Photoresponse of Ge Devices Light from a halogen lamp was passed through a monochromator, modulated by a mechanical chopper, and focused on the surface of the devices via optical fibers. The incident power was determined by illuminating an aperture with a diameter identical to the device size and measuring the power of the light passing through the aperture. The accuracy of this procedure was corroborated using commercial calibrated InGaAs and Ge p-i-n photodiodes. The photocurrent generated in the diode induced a voltage on a load resistor (10 k) that was directly measured by a lock-in amplifier. The voltage and resistance were then used to calculate the photocurrent, which, divided by the power of the incident light, yielded

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10−1 dark

10−2

0.1 8 6

illuminated

10−3 900 nm measured 900 nm theory, ηc = 0.34

2

350 nm measured

10−4 10−5

350 nm theory, ηc = 0.46 0.01 1000

1200

1400 Wavelength (nm)

the responsivity R. The external quantum efficiency (EQE) is obtained as EQE = 1240 nm × R/λ (o nm). Fig. 8 compares the EQE spectral dependence for the p-i-n device with a ∼350-nm intrinsic layer (squares) and the n-i-p device with a ∼900-nm intrinsic layer (circles). The data in both cases were obtained from 300-μm-diameter mesas at zero nominal bias. The sharp decline of both curves above 1600 nm is due to the direct bandgap of the material (1550 nm). Beyond 1600 nm, below the direct bandgap, the signal is reduced to essentially zero, as only indirect transition absorption contributes. At wavelengths shorter than the bandgap, clear oscillations are observed, suggesting interference effects. Since the oxide thickness is the same in both samples, but the oscillation period is different, the Ge layers themselves must play a role in the interference phenomena. To account for these effects we write the EQE as EQE = fηc (1−T+−R+ )+exp(αGe dt op )fηc T+ Rback (1−T−−R− ). (1) Here, ηc is the collection efficiency, T+ (R+ ) is the transmittance (reflectance) of the entire oxide/Ge stack on Si under illumination from the top surface, Rback is the reflectance at the back surface of the Si wafer, and T− (R− ) is the transmittance (reflectance) of the oxide/Ge stack for illumination by the light reflected from the back surface. The factor 1 − exp (−αGe dint ) exp αGe dt op − exp (−αGe dint )

0.35 0.30 0.25 0.20 0.15

10−6

1600

Fig. 8. EQE plots for two Ge/Si diode devices. Circles and squares represent experimental data. The lines are fits based on (1) in which the Ge layer thicknesses are adjusted to match the EQE oscillations and the collection efficiency is adjusted to fit the measured EQE at 1300 nm.

f =

EQE

4

Current (A)

EQE

2

(2)

gives the fraction of incoming the light absorbed in the intrinsic layer of the structure. Here, αGe is the absorption coefficient of Ge, dt op is the thickness of the top doped layer, and dint is the thickness of the intrinsic layer. The corresponding fraction for the light reflected at the back surface is f exp(αGe dt op ). The transmittances and reflectances are calculated using standard transfer-matrix techniques and tabulated optical constants for SiO2 , Ge, and Si. Notice that (1) treats the light traveling through the Ge device structure as fully coherent, but neglects the coherence between the light traveling toward the back surface of the Si wafer and

−6

0

−5

2 4 6 Reverse Bias (V)

−4

−3 −2 Bias (V)

−1

0

1

Fig. 9. I–V curves of the 900-nm Ge photodiode without and under illumination by a 0.6 mW laser diode operating at 1300 nm. The inset shows the EQE dependence on reverse bias at the same conditions.

the light reflected at this surface. The use of the factor f in (1) is an approximation, a more rigorous treatment within the transfer-matrix formalism is discussed by Prentice [17]. The assumption that only light absorbed in the intrinsic layer contributes to the photocurrent slightly underestimates the quantum efficiency, since light absorbed in the top layer will also contribute as long as the electron–hole pairs are generated at distances from the intrinsic layer which are comparable to the minority carrier diffusion length. The theoretical curves in Fig. 8 are fits to the experimental data in which the collection efficiency is treated as an adjustable parameter and the thickness of the intrinsic layer is adjusted to match the observed EQE oscillations. We find ηc = 0.46 for the thinner p-i-n sample and ηc = 0.34 for the thicker sample. The fitted layer thicknesses (87 nm n-Ge and 860 nm i-Ge in the n-i-p diode, and 103 nm p-Ge and 389 nm i-Ge in the p-i-n diode) are within 4% of the thicknesses determined from ellipsometric measurements on the wafer prior to device processing. The discrepancy between theory and experiment at the onset of absorption in the 1500–1600-nm range is not surprising given the strong temperature dependence of the absorption coefficient in this wavelength range [18], the possibility of residual strain effects associated with processing, and the change in the optical constants of Ge under heavy doping, which are neglected in the simulation. The relatively low value of the collection efficiency at zero bias is probably due to residual doping in the intrinsic region. This interpretation implies that the collection efficiency should increase with reverse bias [4]. In Fig. 9, we show I–V curves for the n-i-p detector with an intrinsic region of 900 nm and a diameter of 300 μm. The curves were measured under dark conditions and illumination with a laser emitting 0.6 mW of 1300-nm radiation. The substantial difference between the illuminated and dark current as shown in Fig. 9 (red and black traces, respectively) allows us to extend the EQE measurements to a large reverse bias. At a bias of −0.25 V, the


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10000 450 nm 350 nm 300 nm 250 nm

α (cm−1)

8000

P GeSn Sn

6000 Si

Ge

4000 2000 0

0.60

0.70 0.80 Energy (eV)

0.90

Fig. 10. Plot of absorption coefficients for a set of Ge-Sn samples with 2% Sn, and various thicknesses showing < 0.2% spread of the measured values, demonstrating the reproducibility of Ge-Sn growth.

collection efficiency of the n-i-p device is ηc = 0.7 and reaches full collection at about −6 V. This behavior has been observed in Ge diodes based on similar heterojunctions designs [19]. Similar measurements of a 300-μm diameter n-i-p detector with a 350-nm intrinsic layer show that, at a bias of −0.25 V, full collection efficiency is achieved. E. Fabrication and Photoresponse of Ge–Sn p-i-n Diodes Ge1−y Sn y films with Sn concentration of 2 at% and thickness up to 1000 nm are grown routinely and reproducibly on 4 in Si(100) substrates via reactions of deuterated stannane (SnD4 ) with digermane (Ge2 H6 ) using ultra-high-vacuum chemical vapor deposition, as described elsewhere [7], [20]. The films are produced at a low temperature of 300–350 ºC and pressure of 300 mTorr, with growth rates approaching 15 nm/min. The resultant materials are characterized for composition, structure, and morphology by Rutherford back scattering (RBS), AFM, XTEM, and high-resolution X-ray diffraction (HR-XRD). The ratio of aligned versus random peak heights in the RBS spectra (χmin ), which measures the degree of crystallinity, decreases from 10% at the interface to 5% at the surface. This indicates a reduction in dislocation density across the thickness of the film. The 5% value is close to the practical limit of ∼3% for a perfect Si crystal, suggesting that most of the defects accommodating the lattice mismatch between the film and the substrate are confined at the interface. This is consistent with XTEM observations, which indicate essentially defect-free films. A post-growth RTA step further reduces the concentration of residual threading defects and improves the mosaic spread of the crystal, as evidenced in HR-XRD experiments by a reduction of the FWHM of the (004) rocking curve to 0.275º. In addition, the HR-XRD scans of the (004) and (224) reflections confirm the full relaxation of the material. The AFM scans reveal flat surface morphology of the Ge0.98 Sn0.02 films, with an rms surface roughness of 0.5 nm on a 20 μm × 20 μm area. Extensive optical characterizations of Ge1−y Sn y samples have demonstrated a very strong compositional dependence of the energy gaps, and, therefore, precise stoichiometry control

Si(100)

250 nm 101

104 Counts

107

Fig. 11. XTEM micrograph of the Ge0.98 Sn0.02 structure and corresponding SIMS elemental profile shown in the left and right panels, respectively.

is key to ensuring reproducibility from a device development and performance perspective. Fig. 10 demonstrates a high degree of consistency in intrinsic Ge–Sn synthesis by comparing ellipsometric measurements of the absorption coefficients for randomly selected Ge0.98 Sn0.02 samples grown under the same conditions. The curves are nearly identical within the error of the instrument, and the scattering of the bandgap values obtained by fitting these data is invariably less than 20 meV regardless of thickness. These results confirm that Ge–Sn layers with narrowly specified bandgaps can be grown with near-perfect reproducibility. Another remarkable characteristic of the Ge–Sn layers in Fig. 10 is that the direct band edge is clearly visible as a steep slope in the absorption. This behavior is due to the narrowing of optical transitions attributed to material quality improvements associated primarily with the post-growth RTA treatment. The development of viable in situ n-doping protocols of the Ge–Sn alloys by P and As has also been demonstrated using single-source P(GeH3 )3 and As(GeH3 )3 precursors, to yield active carrier concentrations above 5 × 1019 cm−3 , as measured by both electrical and optical methods. In situ p-doping was achieved using B2 H6 , and the measured hole concentrations were found to be as high as 1 × 1020 cm−3 . Ge0.98 Sn0.02 p-i-n structures were grown on low-resistivity (ρ = 0.02 cm) p-type (B-doped) Si(100) substrates, and were doped in situ using the previously described softchemistry procedures [8]. The thicknesses of the intrinsic Ge– Sn layer and the n-type epilayer (P-doped) were measured by ellipsometry and RBS to be 350 and 100 nm, respectively. XTEM micrographs of these samples showed a periodic array of edge dislocations confined at the Si interface and relatively few threading defects propagating through the bulk of the film (Fig. 11, left). SIMS elemental profiles revealed a homogeneous distribution of the atoms across the entire stack as well as abrupt and well-defined heterointerfaces (Fig. 11, right). The B and P concentrations of the p-substrate and the n-Ge–Sn layer were measured by SIMS to be 4.3 × 1019 and 1 × 1019 cm−3 , respectively. The latter value was independently confirmed by contactless IR ellipsometry. The B and P concentrations of the intrinsic Ge–Sn layer were found to


10

−0.6 V Ea 0.12 eV

0.1 −1.0 −0.5

−7.5

−0.4 V Ea 0.14 eV

1

O

−7.0

−0.8 V Ea 0.10 eV

−8.0

2

−8.5

−0.2 V Ea 0.18 eV

−9.0 0.0 0.5 Bias (V)

1.0

35

36 37 38 1/kT (eV−1)

39

Fig. 12. (Left) Dark current density plots for the Ge0.98 Sn0.02 photodiodes with mesa diameters ranging from 60 to 350 μm and active layer thickness of 350 nm. (Right) Arrhenius plots (natural log of dark current versus 1/T) of devices with 250 μm diameter. Corresponding activation energies indicate a significant contribution to the dark current from recombination generation.

be below the detection limit, indicating that no discernible interdiffusion of the atoms has occurred at the low deposition temperature of 350 °C. Ge–Sn devices were fabricated using previously reported processing protocols, and their electrical properties were characterized using procedures essentially identical to those described above for Ge. Current density versus voltage measurements were conducted, and representative plots for devices sizes ranging from 60 to 250 μm in diameter are show in Fig. 12. The results indicate that the dark current densities at −1 V bias are ∼1 A/cm2, which is 2–3 orders of magnitude higher than those found in the Ge devices. This can be attributed to both the lower bandgap of the Ge–Sn material and to its higher dislocation density relative to Ge. The excess dark currents are consistent with the intermediate value obtained for the diode ideality factor of 1.5, suggesting a significant contribution from recombination–generation processes. In fact, the defect concentration is estimated to be significantly higher (by approximately two orders of magnitude) relative to the defect concentrations in the Ge samples. The nature of the dark currents can be further investigated from their temperature dependence. The Arrhenius plots in Fig. 12 indicate that the activation energies E a for a typical 250-μm device decrease from 0.18 to 0.09 eV with increasing reverse bias from 0.2 to 1 V, respectively. These values are much lower than those obtained from the pure Ge devices. Since E a