**Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors**

Volodymyr Tetyorkin, Andriy Sukach and Andriy Tkachuk

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/52930

### **1. Introduction**

During the last two decades HgCdTe, InSb and InAs infrared (IR) photodiodes have de‐ veloped rapidly for utilization in second generation thermal-imaging systems. Obviously, they are regarded as the most important candidates for development of third generation systems as well. Despite this fact many problems still exist in manufacturing technology as well as in understanding of physical phenomena in materials and photodiodes. As a re‐ sult, threshold parameters of commercially available IR photodiodes are far from the val‐ ues predicted theoretically.

The concept of band gap engineering have found numerous applications in the fabrication IR devices on II-VI and V III-V semiconductors. For instance, the most important advantage of HgCdTe ternary alloy is ability to tune its energy band gap in wide range. The spectral cutoff of HgCdTe photodiodes can be tailored by adjusting the HgCdTe alloy composition over the 1-30 mm range. Further application of this concept in technology of IR detectors is closely connected with development of GaAs/AlGaAs multiple quantum well detectors and InAs/GaInSb type-II superlattice photodiodes.

To implement the concept of defect engineering, grown-in and process-induced defects must be minimized and passivated or eliminated. Defects in narrow-gap semiconductors are easily introduced either intentionally or unintentionally during crystal growth, sample treat‐ ment and device processing. There are also evidences that these defects are electrically ac‐ tive. So, for controlling parameters and characteristics of infrared photodiodes on narrowgap semiconductors through defect engineering, it is essential to understand physical properties of defects, mechanisms of their interaction and temporal evolution. Electronic

© 2012 Tetyorkin et al.; licensee InTech. This is an open access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 Tetyorkin et al.; licensee InTech. This is a paper distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

properties of native defects and foreign impurities in narrow-gap semiconductors have been of great importance for several decades. As a result of intensive investigations, the primary native defects and the mechanisms of their formation were elucidated. Doping effect of dif‐ ferent impurities has been also recognized. This allows to develop effective methods for con‐ trolling the carrier concentration and type of conductivity in intentionally undoped and doped materials. To some extent the carrier lifetime can be controlled by extrinsic doping. However, in many cases electronic states of defects in these semiconductors are still to be unknown and further investigations are needed.

The main objective of this article is to outline the basic properties of point and extended de‐ fects, their effect on physical properties and threshold parameters of infrared photodiodes based on HgCdTe, InAs and InSb narrow-gap semiconductors. This article is divided into two parts. The first part is dedicated to technological steps (crystal growth, thermal anneal‐ ing, junction formation) closely connected with defects forming in materials and devices. In the second part original results are analyzed with emphasize on possible participation of dislocations and point defects in the carrier transport mechanisms and recombination proc‐

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

http://dx.doi.org/10.5772/52930

217

The defect structure of narrow-gap Hg1-xCdxTe (hereinafter – HgCdTe) compounds was in‐ tensively investigated both theoretically and experimentally over the past fifty years. The current status of defect states in these semiconductors are reviewed in numerous papers and monographs (see, e.g., Capper and Garland, 2011; Chu and Sher, 2010). HgCdTe crystalline materials are always grown with large deviation from stoichiometry. The equilibrium exis‐ tence region in HgCdTe (x~0.2) is shown to be completely on the Te-rich side (Schaake, 1985). Thus, the most important type of native point defects in undoped materials are Hg vacancies. Residual impurities, Hg interstitials, dislocations and Te precipitates were also observed in as-grown materials. All these defects can exist in neutral or ionized states. Their important characteristics, such as donor or acceptor type, ionization energy, density, spatial

The ab initio calculation of the formation energies of native point defects in HgCdTe (x~0.2) has been made by Berding with co-authors (Berding, 1994, 1995, 2011). The most reliable cal‐ culations are based on the local-density approximation to the density functional theory. The calculations predict the mercury vacancy and the tellurium antisite TeHg as the dominant de‐ fects in the material grown under tellurium-rich conditions. The concentration of native point defects was calculated as a function of Hg pressure using quasichemical formalism. In the the calculation all defects were assumed to be equilibrated at the temperature 500 0C

cal of LPE growth. In the calculations the Hg vacancy and tellurium antisite TeHg are classi‐ fied as acceptor and donor defects, respectively. The calculation also predicts that the concentration of TeHg is comparable with the mercury vacancies concentration. At the same time, as was pointed by Berding, to date there is no experimental confirmation of the pres‐ ence of the tellurium antisite defects in HgCdTe in so large amounts. Also, the calculated concentration of Hg interstitial was found to be too low to explain experimental data on the self diffusion (Berding, 2011). Experimentally the defect structure in undoped HgCdTe (x~0.2) was investigated by Vydyanath and Schaake. It has been shown that the dominant native defects are doubly ionized acceptors associated with Hg vacancies (Vydyanath, 1981; Schaake, 1985). This result is in accordance with the theoretical prediction. Defects in doped

C temperature is typi‐

**2. Native defects and impurities in HgCdTe, InAs and InSb**

and the defect concentrations were assumed to be frozen in. The 500 0

HgCdTe have been reviewed by Shaw and Capper (Shaw and Capper, 2011).

location and temporal variation were investigated.

esses in the photodiodes.

For controlling properties of semiconductors through defect engineering, it is essential to understand the mechanisms of interaction between point and extended defects (disloca‐ tions), as well as to understand their effect on device characteristics. This task seems to be important since alternative substrates (Si, GaAs, sapphire) are widely used in epitaxial tech‐ nology of IR photodiodes. These substrates are very attractive because they are less expen‐ sive and available in large area wafers. The coupling of the Si substrates with Si read-out integrated circuit allows fabrication of very large focal-plane arrays (FPA). Due to the large lattice mismatch between HgCdTe, InSb and InAs and alternative substrates, photodiodes on their base suffer from the high density of dislocations (typically of the order of 106 cm-2). For instance, these defect densities seriously limit application of HgCdTe epitaxial layers for manufacture of high-performance photodiodes for the LWIR and VLWIR spectral regions. The use of buffer layers, temperature cycling and hydrogen passivation is expected to be useful for reduction of the density of dislocations and weakening their effect on the device performance. However, none of these methods has yet been proven to be practical.

A number of physical properties of HgCdTe, such as direct energy gap, ability to obtain both low and high carrier concentrations, high mobility of electrons and holes, low dielec‐ tric constant and extremely small change of lattice constant with composition, makes it possible to grow high quality layers and heterostructures. As a consequence, high-per‐ formance HgCdTe photodiodes on mead-wavelength, long-wavelength and very longwavelength IR regions (MWIR, LWIR and VLWIR) have been developed. The main drawbacks of HgCdTe are technological disadvantages of this material. The most impor‐ tant is a weak Hg–Te bond, which results in bulk, surface and interface instabilities. Uni‐ formity and yield are still issues especially in photodiodes on the LWIR and VLWIR spectral regions. InSb photodiodes on the MWIR spectral region have comparable per‐ formance with photodiodes made of HgCdTe.

Initially, the IR photodiodes were prepared by diffusion, ion implantation or other techni‐ ques which allow preparation of homojunction structures. In these photodiodes, the concen‐ tration of carriers in the lightly doped 'base' region was strongly controlled. With development of epitaxial techniques, homojunction photodiodes were replaced by hetero‐ junction ones. In a heterojunction photodiode the 'base' region are introduced between a wide-gap substrate and a capping layer with wider band gap. Thus, the influence of surface recombinations on the photodiodes performance is weakened. It seems that the most suc‐ cessful application of the band gap and defect engineering concepts in technology of IR pho‐ todiodes is development of double-layer heterojunction photodiodes.

The main objective of this article is to outline the basic properties of point and extended de‐ fects, their effect on physical properties and threshold parameters of infrared photodiodes based on HgCdTe, InAs and InSb narrow-gap semiconductors. This article is divided into two parts. The first part is dedicated to technological steps (crystal growth, thermal anneal‐ ing, junction formation) closely connected with defects forming in materials and devices. In the second part original results are analyzed with emphasize on possible participation of dislocations and point defects in the carrier transport mechanisms and recombination proc‐ esses in the photodiodes.

#### **2. Native defects and impurities in HgCdTe, InAs and InSb**

properties of native defects and foreign impurities in narrow-gap semiconductors have been of great importance for several decades. As a result of intensive investigations, the primary native defects and the mechanisms of their formation were elucidated. Doping effect of dif‐ ferent impurities has been also recognized. This allows to develop effective methods for con‐ trolling the carrier concentration and type of conductivity in intentionally undoped and doped materials. To some extent the carrier lifetime can be controlled by extrinsic doping. However, in many cases electronic states of defects in these semiconductors are still to be

For controlling properties of semiconductors through defect engineering, it is essential to understand the mechanisms of interaction between point and extended defects (disloca‐ tions), as well as to understand their effect on device characteristics. This task seems to be important since alternative substrates (Si, GaAs, sapphire) are widely used in epitaxial tech‐ nology of IR photodiodes. These substrates are very attractive because they are less expen‐ sive and available in large area wafers. The coupling of the Si substrates with Si read-out integrated circuit allows fabrication of very large focal-plane arrays (FPA). Due to the large lattice mismatch between HgCdTe, InSb and InAs and alternative substrates, photodiodes on their base suffer from the high density of dislocations (typically of the order of 106 cm-2). For instance, these defect densities seriously limit application of HgCdTe epitaxial layers for manufacture of high-performance photodiodes for the LWIR and VLWIR spectral regions. The use of buffer layers, temperature cycling and hydrogen passivation is expected to be useful for reduction of the density of dislocations and weakening their effect on the device

performance. However, none of these methods has yet been proven to be practical.

A number of physical properties of HgCdTe, such as direct energy gap, ability to obtain both low and high carrier concentrations, high mobility of electrons and holes, low dielec‐ tric constant and extremely small change of lattice constant with composition, makes it possible to grow high quality layers and heterostructures. As a consequence, high-per‐ formance HgCdTe photodiodes on mead-wavelength, long-wavelength and very longwavelength IR regions (MWIR, LWIR and VLWIR) have been developed. The main drawbacks of HgCdTe are technological disadvantages of this material. The most impor‐ tant is a weak Hg–Te bond, which results in bulk, surface and interface instabilities. Uni‐ formity and yield are still issues especially in photodiodes on the LWIR and VLWIR spectral regions. InSb photodiodes on the MWIR spectral region have comparable per‐

Initially, the IR photodiodes were prepared by diffusion, ion implantation or other techni‐ ques which allow preparation of homojunction structures. In these photodiodes, the concen‐ tration of carriers in the lightly doped 'base' region was strongly controlled. With development of epitaxial techniques, homojunction photodiodes were replaced by hetero‐ junction ones. In a heterojunction photodiode the 'base' region are introduced between a wide-gap substrate and a capping layer with wider band gap. Thus, the influence of surface recombinations on the photodiodes performance is weakened. It seems that the most suc‐ cessful application of the band gap and defect engineering concepts in technology of IR pho‐

todiodes is development of double-layer heterojunction photodiodes.

unknown and further investigations are needed.

216 Photodiodes - From Fundamentals to Applications

formance with photodiodes made of HgCdTe.

The defect structure of narrow-gap Hg1-xCdxTe (hereinafter – HgCdTe) compounds was in‐ tensively investigated both theoretically and experimentally over the past fifty years. The current status of defect states in these semiconductors are reviewed in numerous papers and monographs (see, e.g., Capper and Garland, 2011; Chu and Sher, 2010). HgCdTe crystalline materials are always grown with large deviation from stoichiometry. The equilibrium exis‐ tence region in HgCdTe (x~0.2) is shown to be completely on the Te-rich side (Schaake, 1985). Thus, the most important type of native point defects in undoped materials are Hg vacancies. Residual impurities, Hg interstitials, dislocations and Te precipitates were also observed in as-grown materials. All these defects can exist in neutral or ionized states. Their important characteristics, such as donor or acceptor type, ionization energy, density, spatial location and temporal variation were investigated.

The ab initio calculation of the formation energies of native point defects in HgCdTe (x~0.2) has been made by Berding with co-authors (Berding, 1994, 1995, 2011). The most reliable cal‐ culations are based on the local-density approximation to the density functional theory. The calculations predict the mercury vacancy and the tellurium antisite TeHg as the dominant de‐ fects in the material grown under tellurium-rich conditions. The concentration of native point defects was calculated as a function of Hg pressure using quasichemical formalism. In the the calculation all defects were assumed to be equilibrated at the temperature 500 0C and the defect concentrations were assumed to be frozen in. The 500 0 C temperature is typi‐ cal of LPE growth. In the calculations the Hg vacancy and tellurium antisite TeHg are classi‐ fied as acceptor and donor defects, respectively. The calculation also predicts that the concentration of TeHg is comparable with the mercury vacancies concentration. At the same time, as was pointed by Berding, to date there is no experimental confirmation of the pres‐ ence of the tellurium antisite defects in HgCdTe in so large amounts. Also, the calculated concentration of Hg interstitial was found to be too low to explain experimental data on the self diffusion (Berding, 2011). Experimentally the defect structure in undoped HgCdTe (x~0.2) was investigated by Vydyanath and Schaake. It has been shown that the dominant native defects are doubly ionized acceptors associated with Hg vacancies (Vydyanath, 1981; Schaake, 1985). This result is in accordance with the theoretical prediction. Defects in doped HgCdTe have been reviewed by Shaw and Capper (Shaw and Capper, 2011).

Native defects in HgCdTe, including dislocations and defect complexes, can act as Shockley-Read-Hall (SRH) centers due to their effect on the carrier lifetime. There is a large literature concerning the links of deep defects to the carrier lifetime in HgCdTe (Capper, 1991; Sher, 1991; Capper, 2011; Cheung, 1985, Chu and Sher, 2010). Clear evidence of SRH centers was provided by deep level transient spectroscopy, admittance spectroscopy, thermally stimulat‐ ed current and optical modulation spectroscopy (Polla, 1981, 1981a, 1981b, 1982; Jones, 1982; Schaake, 1983; Mroczkowski, 1981). The centers located at near midgap seems to be common to p-type Hg1CdTe, where the doping is due to mercury vacancies. Summary of impurity and native defect levels experimentally observed in HgCdTe has been done by Litter et. al. (Litter, 1990). The main results of experimental findings are as follows: (i) shallow acceptorlike levels have activation energies between 2 and 20 meV; deep levels have energies 0.25Eg, 0.5Eg and 0.75Eg above the valence band edge. The concentration of donor-like deep centers reported by Polla and co-authors was ranged from approximately 0.1NA to 10NA, where NA is the shallow acceptor concentration. The values of the cross section for electrons and holes were in the range 10-15-10-16 cm2 and 10-17-10-18 cm2 , respectively. The origin of the SRH cen‐ tres in HgCdTe is still not clear. The vacancy-doped materials with approximately the same carrier concentration, but manufactured at different temperature conditions, may exhibit different lifetimes. At the same time, the correlation has been found between the SRH re‐ combination centre densities and the Hg vacancies concentration (Capper, 2011).

depends on the material composition, concentration of acceptors and degree of compensa‐

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

http://dx.doi.org/10.5772/52930

219

Electronic properties of extended defects in HgCdTe, InSb and InAs are less investigated in comparison with point defects. It is known that II-VI semiconductors are more ionically bonded as III-V covalent semiconductors. As a result, they can be easily plastically de‐ formed at room and lower temperatures (Holt and Yacobi, 2007). The photoplastic effect, discovered by Osipiyan and Savchenko (1968), is conclusive proof that dislocations in II-VI materials are electrically charged and that charge is largely electronically determined. The dislocation core contains broken bonds so the dislocation line may generally be charged negatively. The broken bonds are chemically reactive and electrostatic interaction between the charged dislocation lines and ionized point defects and short-range chemical bonding ef‐ fects may occur. This tends to reduce (neutralize or passivate) the electrical effects of dislo‐ cations especially in the ionically bonded II-VI compounds. The charge states of the dislocation can be altered illumination and other means of carrier injection and this can change dislocation line charges as well as dislocation mobility. Yonenaga (1998) compared the dynamics of dislocations in InAs with those in other semiconductors, including narrowgap and wide-gap II-VI compounds. It has been shown that the activation energies for dislo‐ cation motion depend linearly on the band-gap with an apparent distinction between different types of semiconductors. The activation energy is lower in ionically bonded II-VI semiconductors compare to III-Vs. Thus, dislocations would be expected to be the most mo‐ bile in II-VI narrow-gap semiconductors. The long-range electrostatic interaction strength‐ ens the attraction of dislocations to their Cottrell impurity atmospheres and strengthens the

Dislocations can also act as SRH centers in HgCdTe and III-V narrow-gap semiconductors. The electrical activity of dislocations can be attributed to Cottrell atmosphere, or to dangling bonds in the dislocation core. The ability of the strain fields of dislocations in HgCdTe to capture significant amounts of impurities has been investigated in the early work by Schaake (Schaake, 1983). It has been argued that the origin of the electrical activity is dan‐ gling bond states on the dislocations rather that the impurities. The high Peierls stress ob‐ served in II-VI compound semiconductors supports this argument. This stress has been attributed to the ability of dangling bonds to heal themselves along the dislocation core. The reduced lifetime and mobility caused by dislocations in n- and p-type HgCdTe has been pointed out in several papers (Lopes, 1993; Yamamoto, 1985; Shin, 1991). This reduction has been ascribed to dangling bonds which provide SRH centers. Dislocations also affect the dark currents in LWIR photodiodes operating at low temperatures (<77 K) because they are believed to produce mid-gap states in the band gap (Arias, 1989; Tregilgas, 1988). The de‐ crease of the differential resistance-area product at zero bias voltage, R0A, in the presence of high dislocations densities has been reported by Johnson et al. (Johnson, 1992). As the tem‐ perature decreases the effect of dislocations was found to be more significant. At 77 K the decrease of R0A begins at the dislocation density of the order of 106 cm−2, whereas at 40 K it is affected by the presence of one or more dislocations. The scatter in the R0A data may be associated with the presence of pairs of 'interacting' dislocations, which may be more effec‐

tion. The calculated energy for the shallow acceptors is 11(14) meV for x=0.20 (0.3).

pinning effect opposing dislocation motion.

The behavior of extrinsic defects in HgCdTe is important from several reasons. Manufacture of photodiodes with improved characteristics requires intentionally doped materials with controllable concentrations of acceptors and donors instead of vacancies-doped material. This is caused by several reasons such as instability of the vacancy-doped material and low ability to control the concentration of free carriers. Also, Hg vacancies or complexes with their participation may be responsible not only for the carrier lifetime, but also they can en‐ hance the trap-assisted tunneling, giving rise to excess dark current in infrared photodiodes. Shaw and Capper (Shaw and Capper, 2011) provides a complete summary of the work on dopants in bulk material and epitaxial layers. Indium and iodine are most frequently used as a well-controlled donors for preparation of n-type bulk material and epitaxial layers. Both are incorporated as shallow single donors occupying metal and tellurium lattice sites (In and I, respectively). Indium has moderately high diffusivity whereas diffusivity of iodine is rather low. Group V elements have low diffusivity and hence are ideal as p-type dopants. Approximately 100% activity is found for In and I concentrations up to ~1018 cm−3. Vydya‐ nath (1991) argued that group V elements (P, As, Sb, Bi) are incorporated as donors under Te-rich (Hg-deficient) conditions of growth. Under Hg saturated anneal at 500 0 C there is enough energy to move these elements from metal lattice cites to Te or interstitial sites. The group V impurities if occupy Te sites act as shallow acceptors.

Shallow impurities are known to determine the concentration of free carriers in semiconduc‐ tors. The ionization energy of a hydrogen-like donors in HgCdTe were calculated as a func‐ tion of composition and concentration of defects (Capper, 2011). Due to low effective masses of electrons this energy is too small to be detected experimentally at 77 K (e.g., the ionization energy is 0.30 (0.85) meV for the composition x=0.2 (0.3). The ionization energy for acceptors depends on the material composition, concentration of acceptors and degree of compensa‐ tion. The calculated energy for the shallow acceptors is 11(14) meV for x=0.20 (0.3).

Native defects in HgCdTe, including dislocations and defect complexes, can act as Shockley-Read-Hall (SRH) centers due to their effect on the carrier lifetime. There is a large literature concerning the links of deep defects to the carrier lifetime in HgCdTe (Capper, 1991; Sher, 1991; Capper, 2011; Cheung, 1985, Chu and Sher, 2010). Clear evidence of SRH centers was provided by deep level transient spectroscopy, admittance spectroscopy, thermally stimulat‐ ed current and optical modulation spectroscopy (Polla, 1981, 1981a, 1981b, 1982; Jones, 1982; Schaake, 1983; Mroczkowski, 1981). The centers located at near midgap seems to be common to p-type Hg1CdTe, where the doping is due to mercury vacancies. Summary of impurity and native defect levels experimentally observed in HgCdTe has been done by Litter et. al. (Litter, 1990). The main results of experimental findings are as follows: (i) shallow acceptorlike levels have activation energies between 2 and 20 meV; deep levels have energies 0.25Eg, 0.5Eg and 0.75Eg above the valence band edge. The concentration of donor-like deep centers reported by Polla and co-authors was ranged from approximately 0.1NA to 10NA, where NA is the shallow acceptor concentration. The values of the cross section for electrons and holes

and 10-17-10-18 cm2

combination centre densities and the Hg vacancies concentration (Capper, 2011).

Te-rich (Hg-deficient) conditions of growth. Under Hg saturated anneal at 500 0

group V impurities if occupy Te sites act as shallow acceptors.

enough energy to move these elements from metal lattice cites to Te or interstitial sites. The

Shallow impurities are known to determine the concentration of free carriers in semiconduc‐ tors. The ionization energy of a hydrogen-like donors in HgCdTe were calculated as a func‐ tion of composition and concentration of defects (Capper, 2011). Due to low effective masses of electrons this energy is too small to be detected experimentally at 77 K (e.g., the ionization energy is 0.30 (0.85) meV for the composition x=0.2 (0.3). The ionization energy for acceptors

tres in HgCdTe is still not clear. The vacancy-doped materials with approximately the same carrier concentration, but manufactured at different temperature conditions, may exhibit different lifetimes. At the same time, the correlation has been found between the SRH re‐

The behavior of extrinsic defects in HgCdTe is important from several reasons. Manufacture of photodiodes with improved characteristics requires intentionally doped materials with controllable concentrations of acceptors and donors instead of vacancies-doped material. This is caused by several reasons such as instability of the vacancy-doped material and low ability to control the concentration of free carriers. Also, Hg vacancies or complexes with their participation may be responsible not only for the carrier lifetime, but also they can en‐ hance the trap-assisted tunneling, giving rise to excess dark current in infrared photodiodes. Shaw and Capper (Shaw and Capper, 2011) provides a complete summary of the work on dopants in bulk material and epitaxial layers. Indium and iodine are most frequently used as a well-controlled donors for preparation of n-type bulk material and epitaxial layers. Both are incorporated as shallow single donors occupying metal and tellurium lattice sites (In and I, respectively). Indium has moderately high diffusivity whereas diffusivity of iodine is rather low. Group V elements have low diffusivity and hence are ideal as p-type dopants. Approximately 100% activity is found for In and I concentrations up to ~1018 cm−3. Vydya‐ nath (1991) argued that group V elements (P, As, Sb, Bi) are incorporated as donors under

, respectively. The origin of the SRH cen‐

C there is

were in the range 10-15-10-16 cm2

218 Photodiodes - From Fundamentals to Applications

Electronic properties of extended defects in HgCdTe, InSb and InAs are less investigated in comparison with point defects. It is known that II-VI semiconductors are more ionically bonded as III-V covalent semiconductors. As a result, they can be easily plastically de‐ formed at room and lower temperatures (Holt and Yacobi, 2007). The photoplastic effect, discovered by Osipiyan and Savchenko (1968), is conclusive proof that dislocations in II-VI materials are electrically charged and that charge is largely electronically determined. The dislocation core contains broken bonds so the dislocation line may generally be charged negatively. The broken bonds are chemically reactive and electrostatic interaction between the charged dislocation lines and ionized point defects and short-range chemical bonding ef‐ fects may occur. This tends to reduce (neutralize or passivate) the electrical effects of dislo‐ cations especially in the ionically bonded II-VI compounds. The charge states of the dislocation can be altered illumination and other means of carrier injection and this can change dislocation line charges as well as dislocation mobility. Yonenaga (1998) compared the dynamics of dislocations in InAs with those in other semiconductors, including narrowgap and wide-gap II-VI compounds. It has been shown that the activation energies for dislo‐ cation motion depend linearly on the band-gap with an apparent distinction between different types of semiconductors. The activation energy is lower in ionically bonded II-VI semiconductors compare to III-Vs. Thus, dislocations would be expected to be the most mo‐ bile in II-VI narrow-gap semiconductors. The long-range electrostatic interaction strength‐ ens the attraction of dislocations to their Cottrell impurity atmospheres and strengthens the pinning effect opposing dislocation motion.

Dislocations can also act as SRH centers in HgCdTe and III-V narrow-gap semiconductors. The electrical activity of dislocations can be attributed to Cottrell atmosphere, or to dangling bonds in the dislocation core. The ability of the strain fields of dislocations in HgCdTe to capture significant amounts of impurities has been investigated in the early work by Schaake (Schaake, 1983). It has been argued that the origin of the electrical activity is dan‐ gling bond states on the dislocations rather that the impurities. The high Peierls stress ob‐ served in II-VI compound semiconductors supports this argument. This stress has been attributed to the ability of dangling bonds to heal themselves along the dislocation core. The reduced lifetime and mobility caused by dislocations in n- and p-type HgCdTe has been pointed out in several papers (Lopes, 1993; Yamamoto, 1985; Shin, 1991). This reduction has been ascribed to dangling bonds which provide SRH centers. Dislocations also affect the dark currents in LWIR photodiodes operating at low temperatures (<77 K) because they are believed to produce mid-gap states in the band gap (Arias, 1989; Tregilgas, 1988). The de‐ crease of the differential resistance-area product at zero bias voltage, R0A, in the presence of high dislocations densities has been reported by Johnson et al. (Johnson, 1992). As the tem‐ perature decreases the effect of dislocations was found to be more significant. At 77 K the decrease of R0A begins at the dislocation density of the order of 106 cm−2, whereas at 40 K it is affected by the presence of one or more dislocations. The scatter in the R0A data may be associated with the presence of pairs of 'interacting' dislocations, which may be more effec‐ tive in reducing the R0A than individual dislocations. The excess current in photodiodes caused by dislocations may be the source of 1*/f* noise.

**Impurity, structural**

As divacancy+

As monovacancy+ residual impurity

**defect Type**

neutral

**Ionization energy, meV**

S, Se, Te D EC– (2–3) - Voronina, 1999

Zn A EV+10 - Kesamanly, 1968

Cd A EV+20 - Galkina, 1966 Cd A EV+11 - Iglitsyn, 1968 Mg A - - Voronina, 2004

residual impurity neutral complex - (7 ± 2)·1016 Mahony, J. and Maseher, P., 1977

VIn A1 EV+100 - Mahony, J. and Maseher, P., 1977 VIn A2 EV+130 - Mahony, J. and Maseher, P., 1977 VIn A3 EV+230 - Mahony, J. and Maseher, P., 1977 VAs D EC+0.03Eg - Bynin, M.A. and Matveev, Yu.A., 1985 VAs D EC-Eg/4 - Bynin, M.A. and Matveev, Yu.A., 1985 Cr <sup>D</sup> EC–160 - Omel'yanovskii,1975; Balagurov, 1977;

Mn A EV + (28-30) - Adrianov, 1977; 26 Gheorghitse,1989

Zn A EV +25 - Guseinov, 1969; Guseinov, 1997

Cu D EC-< 7 - Karataev, 1977 Pb N - - Baranov, 1992; Baranov, 1993

Ge (amphoteric) D (Ge sub.In) EC-< 7 - Guseva, 1974; Guseva, 1975 Ge (amphoteric) A (Ge sub. As) EV+14 - Guseva, 1974; Guseva, 1975 Si (amphoteric) D (Si sub. In) EC-< 7 - Guseva, 1974; Guseva, 1975 Si (amphoteric) A (Si sub. As) EV+20 - Guseva, 1974; Guseva, 1975

Structural defect - 110±20 (0.5÷10)·1014 Fomin, 1984 - - 150±20 ≤ 5·10<sup>14</sup> Fomin, 1984 - - ~ 220 (1÷3)·1015 Ilyenkov, 1992 - - ~Eg/2 (77 К) ~ 1017 Kornyushkin, 1996 - D EC–(100-200) - Baranov, 1992 - D EC–(10-20) - Baranov, 1992 - A EV+50 - Baranov, 1992 - D EC–15 - Baranov, 1993 - A EV+50 - Baranov, 1993 - D EC– (20-30) - Voronina, 1999 - D EC–(90-100) - Voronina, 1999 - A EV+10 - Voronina, 1999a - A EV+20 - Voronina, 1999a - A EV+30 - Voronina, 1999a - A EV+65 - Voronina, 1999a - A EV+20 - Zotova, 1975 - A EV+35 - Zotova, 1975 - A EV+35 - Allaberenov, 1970 - A EV+35 - Esina, 1985 Residual impurity D EC–2 - Baranov, 1992 Dislocation (possibly) A EV+(45-50) - Anisimova, 1969

**Table 1.** Parameters of impurities and structural defects in InAs

Be A EV+< 7 - Lin, 1997; Astahov, 1992; Dobbelaere, 1992

**Concentration, cm-3 Reference**

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

Plitnikas,1982; Adomaytis,1984

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221

complex - ~ 1·1017 Mahony, J. and Maseher, P., 1977

The effect of misfit dislocations on dark currents in high temperature MOCVD HgCdTe infrared heterostructure photodiodes has been investigated by Jóźwikowska et al. (Jóźwi‐ kowska, 2004). It was shown that the most effective current transport mechanism at high temperature in HgCdTe heterostructures is the trap-assistant tunneling. In the photodio‐ des operated at 240 K, this mechanism is predominant at bias voltage that not exceeded 0.1 V. The best fit of experimental data with theoretical predictions for the zero bias differ‐ ential resistance versus temperature has been obtained for rather high dislocation densi‐ ties in the volume of HgCdTe layer ~5 107 cm-<sup>2</sup> . To a certain extent electrical activity of dislocation can be reduced by passivation (Boieriu, 2006). It has been showed that incor‐ poration of H in In-doped HgCdTe (x = 0.2) epilayers, through exposure to an electron cy‐ clotron resonance (ECR) H plasma, the lifetime increases by a factor of 10. The increase was attributed to H passivation of the dangling bonds and is only effective for high dislo‐ cation densities (~107 cm−2).

The early studies of native defects and impurities in InSb and InAs have been summarized by Milnes (Milnes, 1973). The energy of shallow impurities in InAs and InSb was calculated with extension of effective mass theory (Baldereschi and Lipari, 1974). Several deep donors of undetermined origin have been observed in InSb. Copper atoms segregated at disloca‐ tions in InAs are apparently electrically inactive, but on heating they diffuse away and can be frozen into the lattice as electrically active centers. They return to the dislocations during the low-temperature anneal. A similar effect was observed with Cu in InSb. A number of undetermined deep defects was found in InSb and InAs (Madelung, 2003).

The nature of intrinsic point defects in InAs single crystals has been studied by several groups (Bublik, 1977, 1979, 1979a; Karataev, 1977; Mahony and Maseher, 1977). In special‐ ly undoped InAs single crystals grown by direct solidification and Czochralski methods precision measurement of density and lattice constant has been made in order to deter‐ mine the type and concentration of interstitial atoms and vacancies depending on the con‐ tent of arsenic in InAs melt. The difference in concentrations of VAs and Asi for InAs was found to be of the order of 3 1018 cm-3. For InAs grown from the melt of stoichiometric composition, this difference does not exceed 1·1017 cm-3. It was also established the effect of point defects on structural and recombination properties of InAs single crystals. It is shown that InAs crystals were n-type conductivity and electron concentration increased from 1.4·1016 to 2.5·1016 cm-3 (T = 77 K) with increasing of As content in the growth melt. It has been concluded that the concentration of electrons is determined by intrinsic defects and complexes composed of native defects (vacancies) and background impurities. The ef‐ fect of annealing on electrical properties of undoped indium arsenide has been investigat‐ ed by Karataev (1977). The annealing was made in the temperature range 300-900 °C for 1-100 h. It was found that the annealing increases the concentration of electrons. For ex‐ ample, for the annealing temperature ~900 °C the concentration of electrons increased from 2·1016 cm-3 to 2·1017 cm-3.


**Table 1.** Parameters of impurities and structural defects in InAs

tive in reducing the R0A than individual dislocations. The excess current in photodiodes

The effect of misfit dislocations on dark currents in high temperature MOCVD HgCdTe infrared heterostructure photodiodes has been investigated by Jóźwikowska et al. (Jóźwi‐ kowska, 2004). It was shown that the most effective current transport mechanism at high temperature in HgCdTe heterostructures is the trap-assistant tunneling. In the photodio‐ des operated at 240 K, this mechanism is predominant at bias voltage that not exceeded 0.1 V. The best fit of experimental data with theoretical predictions for the zero bias differ‐ ential resistance versus temperature has been obtained for rather high dislocation densi‐

dislocation can be reduced by passivation (Boieriu, 2006). It has been showed that incor‐ poration of H in In-doped HgCdTe (x = 0.2) epilayers, through exposure to an electron cy‐ clotron resonance (ECR) H plasma, the lifetime increases by a factor of 10. The increase was attributed to H passivation of the dangling bonds and is only effective for high dislo‐

The early studies of native defects and impurities in InSb and InAs have been summarized by Milnes (Milnes, 1973). The energy of shallow impurities in InAs and InSb was calculated with extension of effective mass theory (Baldereschi and Lipari, 1974). Several deep donors of undetermined origin have been observed in InSb. Copper atoms segregated at disloca‐ tions in InAs are apparently electrically inactive, but on heating they diffuse away and can be frozen into the lattice as electrically active centers. They return to the dislocations during the low-temperature anneal. A similar effect was observed with Cu in InSb. A number of

The nature of intrinsic point defects in InAs single crystals has been studied by several groups (Bublik, 1977, 1979, 1979a; Karataev, 1977; Mahony and Maseher, 1977). In special‐ ly undoped InAs single crystals grown by direct solidification and Czochralski methods precision measurement of density and lattice constant has been made in order to deter‐ mine the type and concentration of interstitial atoms and vacancies depending on the con‐ tent of arsenic in InAs melt. The difference in concentrations of VAs and Asi for InAs was found to be of the order of 3 1018 cm-3. For InAs grown from the melt of stoichiometric composition, this difference does not exceed 1·1017 cm-3. It was also established the effect of point defects on structural and recombination properties of InAs single crystals. It is shown that InAs crystals were n-type conductivity and electron concentration increased from 1.4·1016 to 2.5·1016 cm-3 (T = 77 K) with increasing of As content in the growth melt. It has been concluded that the concentration of electrons is determined by intrinsic defects and complexes composed of native defects (vacancies) and background impurities. The ef‐ fect of annealing on electrical properties of undoped indium arsenide has been investigat‐ ed by Karataev (1977). The annealing was made in the temperature range 300-900 °C for 1-100 h. It was found that the annealing increases the concentration of electrons. For ex‐ ample, for the annealing temperature ~900 °C the concentration of electrons increased

undetermined deep defects was found in InSb and InAs (Madelung, 2003).

. To a certain extent electrical activity of

caused by dislocations may be the source of 1*/f* noise.

220 Photodiodes - From Fundamentals to Applications

ties in the volume of HgCdTe layer ~5 107 cm-<sup>2</sup>

cation densities (~107 cm−2).

from 2·1016 cm-3 to 2·1017 cm-3.


The greatest effect of the anneal was observed in InAs, grown from the melt with an excess

contribute to the effective change in the concentration of electrons at the annealing process. Investigation of defect complexes in InAs using positron annihilation has been presented by Mahony and Maseher (1977). The neutral complexes with the concentration of (7±2)∙10<sup>16</sup> cm-3 has been found. Moreover, these complexes are composed of two vacancies at high temperatures whereas at low temperatures only one vacancy is participated in a complex. Both complexes are stable up to temperatures ~850 ºC. Theoretical calculations of energy levels related with In and As vacancies was also carried out (Mahony and Maseher, 1977). It is shown that In and As vacancies are acceptors and donors, respectively. Growing the high resistivity InAs LPE films doped with chromium and study its electrical and photovoltaic properties was also reported (Omel'yanovskii,1975; Balagurov, 1977; Plitnikas,1982; Ado‐ maytis,1984). Based on investigations of deep centers in epitaxial layers of InAs, the devel‐ opment of high-quality MOS-structures with a density of surface states ≤2 1010 cm-2 eV-1 has been realized. Linear and matrix IR photodetector hybrid assemblies with the temperature

The deviation from the stoichiometry in InSb has been investigated by Abaeva et al. (Abae‐ va, 1987). The authors make a precise determination of the temperature dependence of the InSb lattice constant by Bond's method. It is shown that In and Sb vacancies are the principal intrinsic point defects that determine the deviation from stoichiometry in InSb and reduce the lattice constant. The mid-gap defect states were observed in InSb bulk crystals (Ehem‐ berdyeva, 1982; Shepelyna, 1992). Their energy does not depend on the chemical nature of dopant and the doping level. Defect complex composed of indium and oxygen atoms with energy EC-50 meV has been found by Zaitov et al. (Zaitov, 1981). The low temperature (~200

C) annealing in an atmosphere of inert gas or saturated vapor of lead results in 5-10 times increasing of the carrier lifetime in n-InSb (Tsitsina,1975; Strelnikova, 1993). A model of the conductivity type conversion with participation of fluorine has been proposed (Blaut-Bla‐ chev,1979; Kevorkov, 1980). Some parameters of impurities and deep defects in InAs and

Various methods have been proposed for growing II-VI and III-V semiconductor crystals. These methods can be roughly classified into the melt growth, solution and vapor growth methods. Each of the proposed growth methods has its own advantages and disadvantages. For growing II-VI semiconductor compounds there exist a problem of high dissociation pressure which leads to evaporation of a volatile component and hence deviation from stoi‐ chiometry. The deviation from stoichiometry introduces lattice defects since there is a defi‐ ciency of one component. If there is a large deviation from stoichiometry, vacancies can combine with the residual impurities or with the impurities which are later introduced into

to 105 cm-2 were found to

http://dx.doi.org/10.5772/52930

223

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

of indium or arsenic atoms. Dislocations with densities from 103

resolution of 4 - 8 mK were prepared (Kuryshev, 2001, 2009).

InSb are summarized in Table. 1 and 2.

**3.1. Growth techniques**

**3. Defect-related manufacturing processes**

0

a. in Table 1 and 2 residual impurities and structural defects have unknown identity;

b. doping behavior is indicated by A, acceptor, D, donor; A1,A2 and A3 means singly, doubly and triple ionized acceptor.

**Table 2.** Parameters of impurities and structural defects in InSba,b

The greatest effect of the anneal was observed in InAs, grown from the melt with an excess of indium or arsenic atoms. Dislocations with densities from 103 to 105 cm-2 were found to contribute to the effective change in the concentration of electrons at the annealing process. Investigation of defect complexes in InAs using positron annihilation has been presented by Mahony and Maseher (1977). The neutral complexes with the concentration of (7±2)∙10<sup>16</sup> cm-3 has been found. Moreover, these complexes are composed of two vacancies at high temperatures whereas at low temperatures only one vacancy is participated in a complex. Both complexes are stable up to temperatures ~850 ºC. Theoretical calculations of energy levels related with In and As vacancies was also carried out (Mahony and Maseher, 1977). It is shown that In and As vacancies are acceptors and donors, respectively. Growing the high resistivity InAs LPE films doped with chromium and study its electrical and photovoltaic properties was also reported (Omel'yanovskii,1975; Balagurov, 1977; Plitnikas,1982; Ado‐ maytis,1984). Based on investigations of deep centers in epitaxial layers of InAs, the devel‐ opment of high-quality MOS-structures with a density of surface states ≤2 1010 cm-2 eV-1 has been realized. Linear and matrix IR photodetector hybrid assemblies with the temperature resolution of 4 - 8 mK were prepared (Kuryshev, 2001, 2009).

The deviation from the stoichiometry in InSb has been investigated by Abaeva et al. (Abae‐ va, 1987). The authors make a precise determination of the temperature dependence of the InSb lattice constant by Bond's method. It is shown that In and Sb vacancies are the principal intrinsic point defects that determine the deviation from stoichiometry in InSb and reduce the lattice constant. The mid-gap defect states were observed in InSb bulk crystals (Ehem‐ berdyeva, 1982; Shepelyna, 1992). Their energy does not depend on the chemical nature of dopant and the doping level. Defect complex composed of indium and oxygen atoms with energy EC-50 meV has been found by Zaitov et al. (Zaitov, 1981). The low temperature (~200 0 C) annealing in an atmosphere of inert gas or saturated vapor of lead results in 5-10 times increasing of the carrier lifetime in n-InSb (Tsitsina,1975; Strelnikova, 1993). A model of the conductivity type conversion with participation of fluorine has been proposed (Blaut-Bla‐ chev,1979; Kevorkov, 1980). Some parameters of impurities and deep defects in InAs and InSb are summarized in Table. 1 and 2.

#### **3. Defect-related manufacturing processes**

#### **3.1. Growth techniques**

Impurity,

Ini

structural defect **Type Ionization energy,**

222 Photodiodes - From Fundamentals to Applications

**meV**

**Concentration,**

Zn A EV + 8 - Ismailov, 1969 Zn A EV + 10 - Pehek and Levinstein, 1965 Ge A EV+ (16-19) - Ismailov, 1969 Ag1 A1 EV+27 - Pehek and Levinstein, 1965 Ag2 A2 EV+50 - Pehek and Levinstein, 1965 Au A EV+43 - Pehek and Levinstein, 1965 Cu A EV+ (57-64) - Valyashko, 1975 Cu A EV+(23-27) Kevorkov, 1980 Cu, Ag, Au A EV+67 3·1014 Korotin, 1976 Structural defect D EC–48 - Valyashko, 1975 - A EV+ (7-9) ≤ 6·10<sup>14</sup> Kevorkov, 1980 - A EV+ (5-6) - Nasledov, 1962 - D EC– (7-8) - Nasledov, 1962 - D EC–55 ~ 8·1013 Laff and Fan, 1961 - D EC–60 1011 - 1013 Sipovskaya and Smetannikova, 1984



b. doping behavior is indicated by A, acceptor, D, donor; A1,A2 and A3 means singly, doubly and triple ionized acceptor.

+O2 A EC-50 Zaitov, 1981

a. in Table 1 and 2 residual impurities and structural defects have unknown identity;

**Table 2.** Parameters of impurities and structural defects in InSba,b

**cm-3 Reference**

Various methods have been proposed for growing II-VI and III-V semiconductor crystals. These methods can be roughly classified into the melt growth, solution and vapor growth methods. Each of the proposed growth methods has its own advantages and disadvantages. For growing II-VI semiconductor compounds there exist a problem of high dissociation pressure which leads to evaporation of a volatile component and hence deviation from stoi‐ chiometry. The deviation from stoichiometry introduces lattice defects since there is a defi‐ ciency of one component. If there is a large deviation from stoichiometry, vacancies can combine with the residual impurities or with the impurities which are later introduced into the crystal during a subsequent process. The number of defects is hard to control in this case. Therefore, from the point of perfect crystallography, it is desirable to grow crystals at temperatures as low as possible. This requirement is especially important for HgCdTe al‐ loys, since they are decomposing solids. Also, the liquidus-solidus temperatures of the mer‐ cury telluride and cadmium telluride compounds are different, which causes their segregation as the alloy is frozen from a melt. The resultant variation of the mole fraction of each compounds results in a consequent variation in energy gap as well as electrical and op‐ tical properties throughout the material. In contrast to HgCdTe, the III-V compounds (InSb and InAs) melt congruently, i.e. a liquid and solid having identical compositions are in equi‐ librium at the melting point. Thus, they can be grown directly from the melt, a process com‐ monly used to grow large boules of InAs and InSb.

The LPE is the most matured method for preparation of high-quality HgCdTe epitaxial films (Capper, 2011). Two different approaches have been developed: growth from Te and Hg sol‐

ques have been used to grow both thin and thick films. The tipping and dipping techniques have been implemented by using both tellurium- and mercury-rich solutions, but only tellu‐ rium-rich solutions have been used with the sliding boat. Both Hg- and Te-rich LPE can pro‐ duce material of excellent compositional uniformity and crystalline quality (Capper, 2011). MBE and MOVPE growth of HgCdTe is performed at much lower temperatures compare to

in MBE (Garland, 2011; Maxey, 2011). Due to the weak Hg-Te bond, the evaporation of this material is not congruent. Consequently, Hg re-evaporates from the growth surface, leaving the surface Hg-deficient. Because of a large flux of Hg is necessary in this case, it means that the MBE growth temperature window is extremely small for HgCdTe and precise control of the growth temperature is highly desirable (Arias, 1994; Arias, 1994a). Furthermore, to ob‐ tain a desired cut-off wavelength within ±6%, one must control the Cd composition x within

Routinely produced CdTe and CdZnTe (Zn~ 3-4%) single crystals are used as substrates for MBE and MOVPE growth. Today, the largest commercially available CdZnTe substrates are

more than two 1024x1024 FPAs (Rogalski, 2009). Also, the quality of epitaxial layers is influ‐ enced by poor thermal conductivity, compositional nonuniformity, native defects, surface

The most used dopants for MBE and MOVPE growth are arsenic as the p-type dopant and indium and iodine as the n-type dopants. There are no problems with in-situ incorporation of indium and iodine and low-temperature post-growth annealing is used to optimize the structural and electronic properties of the doped material, but is not required for activation of dopants. The MBE growth of HgCdTe is optimally performed under Te-rich conditions (Arias, 1994), thus Hg vacancies are the dominant native point defects in as-grown layers. Due to this reason the in-situ incorporated As occupy Hg sites in a lattice and annealing is required to its activate as a p-dopant by transferring to Te sites. The standard activation an‐

vacancies. The high-temperature annealing it limits many of the low-temperature-growth advantages of MBE. Therefore, attempts have been made to obtain the p-type activation of

The use of alternative Si and GaAs substrates are attractive in IR FPA technology due to sev‐ eral reasons including available Si substrates with large sizes (up to 300 mm diameter), low‐ er cost and compatibility with semiconductor processing equipment. The match of the coefficient of thermal expansion with Si readout circuit allows fabrication of very large FPAs exhibiting long-term thermal cycle stability. Of course, the large lattice mismatch between HgCdTe and alternative substrates (Si: 19%, GaAs: 14%) causes a high dislocation density

ing of dislocations in the growth direction. The high dislocation density is detrimental to de‐

to mid-107 cm−2) at the epilayer/substrate interface due to propagat‐

C, respectively. Dipping, tipping and sliding boat techni‐

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

area. At this size, the wafers are unable to accommodate

C in MOVPE and around 150-200 0

http://dx.doi.org/10.5772/52930

C under a Hg overpressure to fill the Hg

C

225

utions at 420-600 0

±0.002 (Garland, 2011).

neal is 10 min at ~425 0

(typically from mid-106

limited to approximately 7x7 cm2

C and 380-500 0

LPE. Growth temperature is typically around 200-350 0

roughness and imperfect surface flatness CdZnTe substrates.

C followed by 24h at ~250 0

As either as-grown or after only a low-temperature anneal (Garland, 2011).

Further progress in IR photodiode technology is connected with epitaxial layers. Epitaxial techniques offer the possibility of growing large area layers with good depth and lateral ho‐ mogeneity, abrupt composition and doping profiles, which can be configured to improve the performance of photodiodes. The commonly used methods for preparation of epitaxial films are liquid phase epitaxy (LPE), molecular beam epitaxy (MBE), metal-organic vapor phase epitaxy (MOVPE), and metal-organic chemical vapor epitaxy (MOCVD). LPE method is a near-thermodynamic equilibrium growth technology. It has the advantages of relatively simple process, high utilization rate of the source material, high crystalline quality of the ep‐ itaxial films and fast growing. The weakness of LEP is that it cannot be used for precision controlled growth of very thin films of nano-scale. In other words, it is not applicable to the growth of superlattices or quantum-well devices and other complex micro-structure materi‐ als. In addition, the morphology of materials grown by LPE is usually worse than that grown by MOCVD or MBE. In these techniques epitaxial growth is performed at low tem‐ peratures, which makes it possible to reduce the native defects density. Obviously, for epi‐ layers the choice of substrate is decisive for minimizing effect of misfit dislocations and interdiffusion at the epilayer/substrate heterojunction. This problem is especially important for InSb and InAs because of the lack of lattice-matched III-V wide-gap semiconductors which may be used as substrates. To overcome this difficulty, the use of alternative sub‐ strates (Si, GaAs, supphire) attracts great attention. Successful implementation of epitaxy of narrow-gap semiconductors on Si substrates can directly lead to possibility of realization of multicolor, monolithic focal plane arrays.

#### *3.1.1. Growth of HgCdTe*

Several historical reviews of the development of bulk HgCdTe have been published (Maier and Hesse, 1980; Capper, 1994; Capper, 1997). Bulk growth of HgCdTe is rather difficult due to the high vapor pressure of Hg at the crystal melting point (about 950 0 C), caused by weak Hg-Te bond in a crystal. For bulk material the Bridgman, solid state recrystallization, and travelling heater methods have been developed. The most successful implementation of the bulk growth technique was the travelling heater method which allows to grow up to 5-cm diameter single crystals of high structural and electrical quality. Bulk crystal growth is cur‐ rently used to support the first generation of photoconductive HgCdTe array fabrication.

The LPE is the most matured method for preparation of high-quality HgCdTe epitaxial films (Capper, 2011). Two different approaches have been developed: growth from Te and Hg sol‐ utions at 420-600 0 C and 380-500 0 C, respectively. Dipping, tipping and sliding boat techni‐ ques have been used to grow both thin and thick films. The tipping and dipping techniques have been implemented by using both tellurium- and mercury-rich solutions, but only tellu‐ rium-rich solutions have been used with the sliding boat. Both Hg- and Te-rich LPE can pro‐ duce material of excellent compositional uniformity and crystalline quality (Capper, 2011).

the crystal during a subsequent process. The number of defects is hard to control in this case. Therefore, from the point of perfect crystallography, it is desirable to grow crystals at temperatures as low as possible. This requirement is especially important for HgCdTe al‐ loys, since they are decomposing solids. Also, the liquidus-solidus temperatures of the mer‐ cury telluride and cadmium telluride compounds are different, which causes their segregation as the alloy is frozen from a melt. The resultant variation of the mole fraction of each compounds results in a consequent variation in energy gap as well as electrical and op‐ tical properties throughout the material. In contrast to HgCdTe, the III-V compounds (InSb and InAs) melt congruently, i.e. a liquid and solid having identical compositions are in equi‐ librium at the melting point. Thus, they can be grown directly from the melt, a process com‐

Further progress in IR photodiode technology is connected with epitaxial layers. Epitaxial techniques offer the possibility of growing large area layers with good depth and lateral ho‐ mogeneity, abrupt composition and doping profiles, which can be configured to improve the performance of photodiodes. The commonly used methods for preparation of epitaxial films are liquid phase epitaxy (LPE), molecular beam epitaxy (MBE), metal-organic vapor phase epitaxy (MOVPE), and metal-organic chemical vapor epitaxy (MOCVD). LPE method is a near-thermodynamic equilibrium growth technology. It has the advantages of relatively simple process, high utilization rate of the source material, high crystalline quality of the ep‐ itaxial films and fast growing. The weakness of LEP is that it cannot be used for precision controlled growth of very thin films of nano-scale. In other words, it is not applicable to the growth of superlattices or quantum-well devices and other complex micro-structure materi‐ als. In addition, the morphology of materials grown by LPE is usually worse than that grown by MOCVD or MBE. In these techniques epitaxial growth is performed at low tem‐ peratures, which makes it possible to reduce the native defects density. Obviously, for epi‐ layers the choice of substrate is decisive for minimizing effect of misfit dislocations and interdiffusion at the epilayer/substrate heterojunction. This problem is especially important for InSb and InAs because of the lack of lattice-matched III-V wide-gap semiconductors which may be used as substrates. To overcome this difficulty, the use of alternative sub‐ strates (Si, GaAs, supphire) attracts great attention. Successful implementation of epitaxy of narrow-gap semiconductors on Si substrates can directly lead to possibility of realization of

Several historical reviews of the development of bulk HgCdTe have been published (Maier and Hesse, 1980; Capper, 1994; Capper, 1997). Bulk growth of HgCdTe is rather difficult due

Hg-Te bond in a crystal. For bulk material the Bridgman, solid state recrystallization, and travelling heater methods have been developed. The most successful implementation of the bulk growth technique was the travelling heater method which allows to grow up to 5-cm diameter single crystals of high structural and electrical quality. Bulk crystal growth is cur‐ rently used to support the first generation of photoconductive HgCdTe array fabrication.

C), caused by weak

to the high vapor pressure of Hg at the crystal melting point (about 950 0

monly used to grow large boules of InAs and InSb.

224 Photodiodes - From Fundamentals to Applications

multicolor, monolithic focal plane arrays.

*3.1.1. Growth of HgCdTe*

MBE and MOVPE growth of HgCdTe is performed at much lower temperatures compare to LPE. Growth temperature is typically around 200-350 0 C in MOVPE and around 150-200 0 C in MBE (Garland, 2011; Maxey, 2011). Due to the weak Hg-Te bond, the evaporation of this material is not congruent. Consequently, Hg re-evaporates from the growth surface, leaving the surface Hg-deficient. Because of a large flux of Hg is necessary in this case, it means that the MBE growth temperature window is extremely small for HgCdTe and precise control of the growth temperature is highly desirable (Arias, 1994; Arias, 1994a). Furthermore, to ob‐ tain a desired cut-off wavelength within ±6%, one must control the Cd composition x within ±0.002 (Garland, 2011).

Routinely produced CdTe and CdZnTe (Zn~ 3-4%) single crystals are used as substrates for MBE and MOVPE growth. Today, the largest commercially available CdZnTe substrates are limited to approximately 7x7 cm2 area. At this size, the wafers are unable to accommodate more than two 1024x1024 FPAs (Rogalski, 2009). Also, the quality of epitaxial layers is influ‐ enced by poor thermal conductivity, compositional nonuniformity, native defects, surface roughness and imperfect surface flatness CdZnTe substrates.

The most used dopants for MBE and MOVPE growth are arsenic as the p-type dopant and indium and iodine as the n-type dopants. There are no problems with in-situ incorporation of indium and iodine and low-temperature post-growth annealing is used to optimize the structural and electronic properties of the doped material, but is not required for activation of dopants. The MBE growth of HgCdTe is optimally performed under Te-rich conditions (Arias, 1994), thus Hg vacancies are the dominant native point defects in as-grown layers. Due to this reason the in-situ incorporated As occupy Hg sites in a lattice and annealing is required to its activate as a p-dopant by transferring to Te sites. The standard activation an‐ neal is 10 min at ~425 0 C followed by 24h at ~250 0 C under a Hg overpressure to fill the Hg vacancies. The high-temperature annealing it limits many of the low-temperature-growth advantages of MBE. Therefore, attempts have been made to obtain the p-type activation of As either as-grown or after only a low-temperature anneal (Garland, 2011).

The use of alternative Si and GaAs substrates are attractive in IR FPA technology due to sev‐ eral reasons including available Si substrates with large sizes (up to 300 mm diameter), low‐ er cost and compatibility with semiconductor processing equipment. The match of the coefficient of thermal expansion with Si readout circuit allows fabrication of very large FPAs exhibiting long-term thermal cycle stability. Of course, the large lattice mismatch between HgCdTe and alternative substrates (Si: 19%, GaAs: 14%) causes a high dislocation density (typically from mid-106 to mid-107 cm−2) at the epilayer/substrate interface due to propagat‐ ing of dislocations in the growth direction. The high dislocation density is detrimental to de‐ vice performance because of the effect of dislocations on minority carrier lifetimes (Shin, 1992). For comparison, in epitaxial layers grown on CdTe and CdZnTe substrates the dislo‐ cation density is less than mid-105 cm-2. The dislocation density below this value is believed to be not a serious problem unless they form clusters under device contacts. Despite these difficulties, epitaxial growth of high quality HgCdTe on 4-inch CdTe/Si substrates has been demonstrated for MWIR applications (Maranowski, 2001). Also, large area high quality HgCdTe epilayers were grown by MBE on 100 mm diameter (211)Si substrates with a CdTe/ ZnTe buffer layer (Bornfreund, 2007). Epilayers of HgCdTe with extremely uniform compo‐ sition and extremely low defects density were demonstrated by Peterson et al. (Peterson, 2006) on 4- and 6-inch diameter silicon substrates.

these films reduction of the concentration of structural defects was also observed. The simul‐ taneous introduction of controlled amounts of lead and rare earth elements makes it possi‐ ble to prepare high resistivity films with the compensation degree 0.6-0.9. In the compensated films the electron concentration 3·1015 cm-3 was achieved. The defect complexes

Various efforts have been made to adopt LPE growth of InSb (Kumagava, 1973; Mengai‐ lis, 1966; Holmes, 1980). However, there is a small number of reports on successful growth of InSb films by this method. Recently Dixit and co-authors reported growth of high-quality films on (001) semi-insulating GaAs substrate in a boat-slider type LPE unit

Heteroepitaxy of InSb and InAs has been achieved on Si and GaAs substrates with MBE and MOCVD technology (Razeghi, 2003). To overcome the lattice mismatch (>19%), the MBE of InSb on Si substrates was performed using CaF2 and stacked BaF2/CaF2 buffer layers. The

μm-thick film grown on a Si substrate with 0.3 μm CaF2 buffer layer. The 77 K mobilities were at least an order of magnitude lower than the room temperature values. This behavior of electron mobility is attributed to electron scattering on dislocations arising from both lat‐

Epitaxial layers of InSb were grown directly on InSb, GaAs and GaAs-coated Si substrates with MBE and a low pressure MOCVD techniques (Razeghi, 2003). The quality of epitax‐ ial films has been shown to depend critically on the growth conditions and preparation of substrates. In order to get high crystal quality InSb and GaAs substrates, directed 20 off the (100) toward (110) direction were used. The X-ray rocking curve FWHM, electron con‐ centration and mobility was found to depend on the thickness of films due to influence of highly dislocated interface. In 3.6 *μ*m thick InSb film the electron mobility was 56000

tion at 77 K was of the order of 1016 cm-3. Excellent uniformity (within the *±*3 arcs varia‐ tion of FWHM) was detected for a 10 μm thick InSb layer grown by MBE on a 3-inch semi-insulating GaAs substrate. The FWHM decreases with thickness as the dislocation density decreases due to the greater distance between the surface and the highly dislocat‐

the films with thickness more than 2 μm. The temperature dependence of the electron mo‐ bility was peaked at 77 K and decreased at lower temperatures due to the dislocation scat‐

The MBE growth of InAs has been reported by several groups (Yano, 1997; Kalem, 1998). InAs epitaxial layers with thicknesses ranging from 0.5 up to 6.2 μm was grown on (100) ori‐ ented semi-insulating GaAs substrates. As in the case of InSb films, the properties of InAs films was influenced by the growth conditions and InAs/GaAs interface structure. The elec‐

cm2

with thickness of 6.2 μm. It is shown that the temperature dependence as well as the magni‐ tude of the mobility can be explained by a combined impurity-phonon-dislocation scattering

ed interface. The 300 K mobility close to that of bulk InSb (75000 cm<sup>2</sup>

/V s (n ≈2 1016 cm-3

)

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

/V·s at 77 K. The background electron concentra‐

/V·s (n=6.1 1015 cm−3) and peaks at about liq‐

/V·s (n=3.1 1015 cm−3) for a InAs layer

was obtained in an 8

http://dx.doi.org/10.5772/52930

227

/V·s) was achieved in

of n-type were also observed in these films.

room temperature electron mobility of 65000 cm<sup>2</sup>

/V·s at 300 K and close to 80000 cm<sup>2</sup>

tering at the InSb/GaAs interface (Razeghi, 2003).

tron mobility at room temperature is 1.8 10<sup>4</sup>

uid-nitrogen temperature with a value of 5.173 10<sup>4</sup> cm2

(Dixit, 2002; Dixit, 2002a).

tice and thermal strains (Liu, 1997).

cm2

The lattice matched and alternative substrates with (100) and (211) orientations are com‐ monly used. The best crystalline quality is obtained on the substrates with the slightly mi‐ soriented surfaces. For instance, the MOVPE growth on (100) substrates misoriented from the (100) plane by a few degrees is useful to suppress the formation of pyramid-shaped mac‐ rodefects, known as hillocks (Maxey, 2011). Today almost all growth is carried out on (211) substrates, which have (111) terraces with (100) steps (Garland and Sporken, 2011).

The MBE and MOVPE technology of HgCdTe has developed to rather high level at which epitaxial layers grown on bulk CdTe and CdZnTe substrates have characteristics compara‐ ble to those of LPE material.

#### *3.1.2. Growth of InSb and InAs*

Due to the relatively low melting point of (525 0 C) and small saturation vapor pressure, InSb bulk crystals can be grown using different growth techniques: Czochralski and horizontal Bridgman technique, travelling heater method and zone melting (Hulme and Mullin, 1962; Liang, 1966; Parker, 1965; Benz and Müller, 1975; Bagai, 1983). Technology of InSb bulk crys‐ tals are well matured. InSb is the most perfect material among the III-V semiconductors available to data. Typically 10-100 cm-2 etch pit density is specified in a commercially availa‐ ble InSb. However, the best result is only 1 etch pit in 50 cm-2. In the ultra-high pure InSb bulk crystals the carrier concentration can be lower 1013 cm-3.

In contrast to indium antimonide, InAs possesses an appreciable vapor pressure at the melt‐ ing point (943 0 C), because of the equilibrium vapor is constituted of the more volatile com‐ ponent (As) whose condensation and sublimation temperature lies below the melting point. However, even at this case, there is a strong tendency to form stoichiometric compound. InAs bulk crystals can be grown using liquid covering Czochralski or vertical gradient freeze method. Because of purification of InAs is more difficult than InSb, the residual elec‐ tron concentration of InAs bulk crystals is about of 2·1016 cm-3.

Gettering effect of lead and rare earth elements (ytterbium and gadolinium) in LPE of InAs has been studied by several groups ( Baranov, 1992, 1993; Voronina, 1999; Gao, 1999). The gettering effect of lead is attributed to the formation of stable insoluble aggregates com‐ posed of indium tellurides, selenides and sulphides. As a result, epitaxial films with the con‐ centration of electrons ∼<sup>15</sup> cm-3 and mobility 9.1 104 cm2 /V s at 77 K have been grown. In these films reduction of the concentration of structural defects was also observed. The simul‐ taneous introduction of controlled amounts of lead and rare earth elements makes it possi‐ ble to prepare high resistivity films with the compensation degree 0.6-0.9. In the compensated films the electron concentration 3·1015 cm-3 was achieved. The defect complexes of n-type were also observed in these films.

vice performance because of the effect of dislocations on minority carrier lifetimes (Shin, 1992). For comparison, in epitaxial layers grown on CdTe and CdZnTe substrates the dislo‐ cation density is less than mid-105 cm-2. The dislocation density below this value is believed to be not a serious problem unless they form clusters under device contacts. Despite these difficulties, epitaxial growth of high quality HgCdTe on 4-inch CdTe/Si substrates has been demonstrated for MWIR applications (Maranowski, 2001). Also, large area high quality HgCdTe epilayers were grown by MBE on 100 mm diameter (211)Si substrates with a CdTe/ ZnTe buffer layer (Bornfreund, 2007). Epilayers of HgCdTe with extremely uniform compo‐ sition and extremely low defects density were demonstrated by Peterson et al. (Peterson,

The lattice matched and alternative substrates with (100) and (211) orientations are com‐ monly used. The best crystalline quality is obtained on the substrates with the slightly mi‐ soriented surfaces. For instance, the MOVPE growth on (100) substrates misoriented from the (100) plane by a few degrees is useful to suppress the formation of pyramid-shaped mac‐ rodefects, known as hillocks (Maxey, 2011). Today almost all growth is carried out on (211)

The MBE and MOVPE technology of HgCdTe has developed to rather high level at which epitaxial layers grown on bulk CdTe and CdZnTe substrates have characteristics compara‐

bulk crystals can be grown using different growth techniques: Czochralski and horizontal Bridgman technique, travelling heater method and zone melting (Hulme and Mullin, 1962; Liang, 1966; Parker, 1965; Benz and Müller, 1975; Bagai, 1983). Technology of InSb bulk crys‐ tals are well matured. InSb is the most perfect material among the III-V semiconductors available to data. Typically 10-100 cm-2 etch pit density is specified in a commercially availa‐ ble InSb. However, the best result is only 1 etch pit in 50 cm-2. In the ultra-high pure InSb

In contrast to indium antimonide, InAs possesses an appreciable vapor pressure at the melt‐

ponent (As) whose condensation and sublimation temperature lies below the melting point. However, even at this case, there is a strong tendency to form stoichiometric compound. InAs bulk crystals can be grown using liquid covering Czochralski or vertical gradient freeze method. Because of purification of InAs is more difficult than InSb, the residual elec‐

Gettering effect of lead and rare earth elements (ytterbium and gadolinium) in LPE of InAs has been studied by several groups ( Baranov, 1992, 1993; Voronina, 1999; Gao, 1999). The gettering effect of lead is attributed to the formation of stable insoluble aggregates com‐ posed of indium tellurides, selenides and sulphides. As a result, epitaxial films with the con‐

C), because of the equilibrium vapor is constituted of the more volatile com‐

C) and small saturation vapor pressure, InSb

/V s at 77 K have been grown. In

substrates, which have (111) terraces with (100) steps (Garland and Sporken, 2011).

2006) on 4- and 6-inch diameter silicon substrates.

226 Photodiodes - From Fundamentals to Applications

Due to the relatively low melting point of (525 0

bulk crystals the carrier concentration can be lower 1013 cm-3.

tron concentration of InAs bulk crystals is about of 2·1016 cm-3.

centration of electrons ∼<sup>15</sup> cm-3 and mobility 9.1 104 cm2

ble to those of LPE material.

*3.1.2. Growth of InSb and InAs*

ing point (943 0

Various efforts have been made to adopt LPE growth of InSb (Kumagava, 1973; Mengai‐ lis, 1966; Holmes, 1980). However, there is a small number of reports on successful growth of InSb films by this method. Recently Dixit and co-authors reported growth of high-quality films on (001) semi-insulating GaAs substrate in a boat-slider type LPE unit (Dixit, 2002; Dixit, 2002a).

Heteroepitaxy of InSb and InAs has been achieved on Si and GaAs substrates with MBE and MOCVD technology (Razeghi, 2003). To overcome the lattice mismatch (>19%), the MBE of InSb on Si substrates was performed using CaF2 and stacked BaF2/CaF2 buffer layers. The room temperature electron mobility of 65000 cm<sup>2</sup> /V s (n ≈2 1016 cm-3 ) was obtained in an 8 μm-thick film grown on a Si substrate with 0.3 μm CaF2 buffer layer. The 77 K mobilities were at least an order of magnitude lower than the room temperature values. This behavior of electron mobility is attributed to electron scattering on dislocations arising from both lat‐ tice and thermal strains (Liu, 1997).

Epitaxial layers of InSb were grown directly on InSb, GaAs and GaAs-coated Si substrates with MBE and a low pressure MOCVD techniques (Razeghi, 2003). The quality of epitax‐ ial films has been shown to depend critically on the growth conditions and preparation of substrates. In order to get high crystal quality InSb and GaAs substrates, directed 20 off the (100) toward (110) direction were used. The X-ray rocking curve FWHM, electron con‐ centration and mobility was found to depend on the thickness of films due to influence of highly dislocated interface. In 3.6 *μ*m thick InSb film the electron mobility was 56000 cm2 /V·s at 300 K and close to 80000 cm<sup>2</sup> /V·s at 77 K. The background electron concentra‐ tion at 77 K was of the order of 1016 cm-3. Excellent uniformity (within the *±*3 arcs varia‐ tion of FWHM) was detected for a 10 μm thick InSb layer grown by MBE on a 3-inch semi-insulating GaAs substrate. The FWHM decreases with thickness as the dislocation density decreases due to the greater distance between the surface and the highly dislocat‐ ed interface. The 300 K mobility close to that of bulk InSb (75000 cm<sup>2</sup> /V·s) was achieved in the films with thickness more than 2 μm. The temperature dependence of the electron mo‐ bility was peaked at 77 K and decreased at lower temperatures due to the dislocation scat‐ tering at the InSb/GaAs interface (Razeghi, 2003).

The MBE growth of InAs has been reported by several groups (Yano, 1997; Kalem, 1998). InAs epitaxial layers with thicknesses ranging from 0.5 up to 6.2 μm was grown on (100) ori‐ ented semi-insulating GaAs substrates. As in the case of InSb films, the properties of InAs films was influenced by the growth conditions and InAs/GaAs interface structure. The elec‐ tron mobility at room temperature is 1.8 10<sup>4</sup> cm2 /V·s (n=6.1 1015 cm−3) and peaks at about liq‐ uid-nitrogen temperature with a value of 5.173 10<sup>4</sup> cm2 /V·s (n=3.1 1015 cm−3) for a InAs layer with thickness of 6.2 μm. It is shown that the temperature dependence as well as the magni‐ tude of the mobility can be explained by a combined impurity-phonon-dislocation scattering mechanism. The dislocation densities of the order of 106 cm−2 were found. High-quality InAs epilayers on the GaAs substrates have been grown by MBE (Chen,2000; Cai, 2003). The growth was carried out as a two-step process: InAs layers were grown under As-rich condi‐ tions on InAs prelayers grown directly on the GaAs substrates under In-rich conditions. The optimized growth condition for this method from the Raman spectroscopy and the low-tem‐ perature photoluminescence was the following: first InAs is grown 20 nm thick under Inrich conditions at 500 0 C with the appropriate V/III ratio of 8, then InAs is continuously grown under As-rich conditions at 500 0 C with the appropriate V/III ratio of 10–23. Also, a two-step growth method consisting of a 400 0 C prelayer followed by deposition of the thick bulk layer at higher growth temperatures has been reported by Watkins et al. (Watkins, 1995). High purity InAs epilayers were grown on GaAs substrates by MOCVD technique. Temperature dependent Hall measurements between 1.8 and 293 K showed a competition between bulk and surface conduction, with average Hall mobilities of 1.2·10<sup>5</sup> cm<sup>2</sup> /V·s at 50 K. Large changes in the temperature dependent transport data are observed several hours after Hall contact formation and appear to be due to passivation of the surface accumulation lay‐ er by native oxide formation. The highest electron mobilities were observed in InAs films grown by MOCVD at reduced growth temperature (Partin, 1991). Electron mobilities as high as 21000 cm2 /V·s at 300 K were obtained for a film only 3.4 μm thick. From the depth dependence of transport properties it has been found that in the grown films electrons are accumulated near the air interface of the film, presumably by positive ions in the native ox‐ ide. The scattering from dislocations was greatly reduced in the surface accumulation layer due to screening by a high density of electrons. These dislocations arise from lattice mis‐ match and interface disorder at the film-substrate interface, preventing these films from ob‐ taining mobility values of bulk indium arsenide. More or less successful attempts to grow quality InAs epitaxial films were also reported in numerous papers (Fukui, 1979; Haywood, 1990; Egan, 1995; Huang, 1995; von Eichel-Streiber, 1997; Watkins, 1997).

time ion implantation is commonly used technique for formation n-on-p and p-on-n homo-

into a vacancy doped p–type material, but B and Be are the most frequently used for this purpose. The doses of 1012–1015 cm–2 and energies of 30–200 keV are used for the junction formation. First devices were fabricated on bulk p-HgCdTe single crystals with the hole con‐ centration of the order of 1016 cm–3. But in the early 1990s bulk crystals were replaced by LPE material. To-day, epitaxial films grown by LPE on CdZnTe lattice matched substrates are the best structural quality material, and they are successfully used in the homojunction technol‐

gap material). In both structures the lightly doped 'base' region with the doping concentra‐ tion below 1016 cm-3 determines the dark current and photocurrent. Indium is most frequent‐ ly used as a well-controlled dopant for n-type doping due to its high solubility and moderately high diffusion. Arsenic proved to be the most successful p-type dopant that are used for fabrication of stable junctions due to very low diffusivity. The important advantage of P+-n structure is that the thermal generation of carriers is effectively reduced in wider

The significant step in the development of HgCdTe photodiodes has been made by Arias with co-authors (Arias, 1993). They proposed the double-layer planar heterostructure (DLPH) photodiodes. The photodiodes were realized by incorporating a buried narrowbandgap active layer in the DLPH configuration. The planar devices were formed using a Hg1–yCdyTe/Hg1–xCdxTe (y>x) heterostructure grown by MBE on CdZnTe substrate. An im‐ portant feature of the DLPH approach is a planar p-doped/n-doped device geometry that in‐ cludes a wide-bandgap cap layer over a narrow-bandgap 'base' layer. The formation of planar photodiodes was achieved by selective implanting of arsenic through a ZnS mask fol‐ lowed by diffusing the arsenic (by annealing at high temperature) through the cap layer into the narrow gap base layer. After that the structures were annealed under Hg overpressure. The first high-temperature annealing was carried out to diffuse the arsenic into the base lay‐ er and to make the doped region p-type by substitution of arsenic atoms on the Te sub-lat‐ tice, while the second low-temperature one was carried out to annihilate Hg vacancies formed in the HgCdTe lattice during high-temperature process. In DLPH photodiodes sig‐ nificant reduction in tunneling current and surface generation-recombination current has been achieved. The architecture of mesa and planar heterostructure photodiodes are shown in Fig.1. The back-illuminated heterostructure photodiodes prepared from both LPE and MBE material grown on CdZnTe substrates have the highest performance achieved to-day

The serious disadvantage of CdZnTe substrates is the thermal expansion coefficient mis‐ match with Si used for the read-out integrated circuit. To overcome this problem, growth of HgCdTe on alternate substrates such as sapphire, Si and GaAs has been developed. LPE, MBE and MOCVD techniques were used for this purpose. However, these alternate substrate ma‐ terials suffer from a large lattice mismatch with HgCdTe, leading to a higher defect density in the HgCdTe material and consequently reducing the detector performance (Bajaj, 2000). De‐ spite these disadvantages the large, 1024-1024 and 2048-2048, HgCdTe FPAs operating in


Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors


http://dx.doi.org/10.5772/52930

229

and heterojunctions. The n+

bandgap material.

(Arias, 1993; Bajaj, 2000; Rogalski, 2000).

ogy. An ion implantation was also adapted to produce P+

Technology of InSb bulk crystals is more mature than HgCdTe. Good quality substrates with more than 7 cm diameter are commercially available (Micklethwaite, 2000). However, technology of InSb and InAs epitaxial layers is not matured to the level suitable for device applications.

#### **3.2. Junction formation techniques**

#### *3.2.1. HgCdTe photodiodes*

Different HgCdTe photodiode structures have been developed including mesa, planar and lateral homojunction and heterojunction structures. The history and current status of HgCdTe IR photodiodes has been reviewed in numerous monographs and review articles (Capper and Elliot, 2001; Norton,1999; Rogalski, 2011; Rogalski et al., 2000; Chu and Sher, 2010; Reine, 2000; Baker and Maxey, 2001; Sher, 1991). The junctions have been formed by numerous techniques including impurity diffusion, ion implantation, growth of doped epi‐ taxial layers from vapor or liquid phase. In the early stages of device technology Hg in- and out-diffusion has been used for the junction formation in HgCdTe bulk material. At present time ion implantation is commonly used technique for formation n-on-p and p-on-n homoand heterojunctions. The n+ -p homojunction can be formed by implantation of different ions into a vacancy doped p–type material, but B and Be are the most frequently used for this purpose. The doses of 1012–1015 cm–2 and energies of 30–200 keV are used for the junction formation. First devices were fabricated on bulk p-HgCdTe single crystals with the hole con‐ centration of the order of 1016 cm–3. But in the early 1990s bulk crystals were replaced by LPE material. To-day, epitaxial films grown by LPE on CdZnTe lattice matched substrates are the best structural quality material, and they are successfully used in the homojunction technol‐ ogy. An ion implantation was also adapted to produce P+ -n heterojunctions (P means wider gap material). In both structures the lightly doped 'base' region with the doping concentra‐ tion below 1016 cm-3 determines the dark current and photocurrent. Indium is most frequent‐ ly used as a well-controlled dopant for n-type doping due to its high solubility and moderately high diffusion. Arsenic proved to be the most successful p-type dopant that are used for fabrication of stable junctions due to very low diffusivity. The important advantage of P+-n structure is that the thermal generation of carriers is effectively reduced in wider bandgap material.

mechanism. The dislocation densities of the order of 106

rich conditions at 500 0

high as 21000 cm2

applications.

**3.2. Junction formation techniques**

*3.2.1. HgCdTe photodiodes*

grown under As-rich conditions at 500 0

228 Photodiodes - From Fundamentals to Applications

two-step growth method consisting of a 400 0

epilayers on the GaAs substrates have been grown by MBE (Chen,2000; Cai, 2003). The growth was carried out as a two-step process: InAs layers were grown under As-rich condi‐ tions on InAs prelayers grown directly on the GaAs substrates under In-rich conditions. The optimized growth condition for this method from the Raman spectroscopy and the low-tem‐ perature photoluminescence was the following: first InAs is grown 20 nm thick under In-

bulk layer at higher growth temperatures has been reported by Watkins et al. (Watkins, 1995). High purity InAs epilayers were grown on GaAs substrates by MOCVD technique. Temperature dependent Hall measurements between 1.8 and 293 K showed a competition

Large changes in the temperature dependent transport data are observed several hours after Hall contact formation and appear to be due to passivation of the surface accumulation lay‐ er by native oxide formation. The highest electron mobilities were observed in InAs films grown by MOCVD at reduced growth temperature (Partin, 1991). Electron mobilities as

dependence of transport properties it has been found that in the grown films electrons are accumulated near the air interface of the film, presumably by positive ions in the native ox‐ ide. The scattering from dislocations was greatly reduced in the surface accumulation layer due to screening by a high density of electrons. These dislocations arise from lattice mis‐ match and interface disorder at the film-substrate interface, preventing these films from ob‐ taining mobility values of bulk indium arsenide. More or less successful attempts to grow quality InAs epitaxial films were also reported in numerous papers (Fukui, 1979; Haywood,

Technology of InSb bulk crystals is more mature than HgCdTe. Good quality substrates with more than 7 cm diameter are commercially available (Micklethwaite, 2000). However, technology of InSb and InAs epitaxial layers is not matured to the level suitable for device

Different HgCdTe photodiode structures have been developed including mesa, planar and lateral homojunction and heterojunction structures. The history and current status of HgCdTe IR photodiodes has been reviewed in numerous monographs and review articles (Capper and Elliot, 2001; Norton,1999; Rogalski, 2011; Rogalski et al., 2000; Chu and Sher, 2010; Reine, 2000; Baker and Maxey, 2001; Sher, 1991). The junctions have been formed by numerous techniques including impurity diffusion, ion implantation, growth of doped epi‐ taxial layers from vapor or liquid phase. In the early stages of device technology Hg in- and out-diffusion has been used for the junction formation in HgCdTe bulk material. At present

between bulk and surface conduction, with average Hall mobilities of 1.2·10<sup>5</sup>

1990; Egan, 1995; Huang, 1995; von Eichel-Streiber, 1997; Watkins, 1997).

C with the appropriate V/III ratio of 8, then InAs is continuously

/V·s at 300 K were obtained for a film only 3.4 μm thick. From the depth

C with the appropriate V/III ratio of 10–23. Also, a

C prelayer followed by deposition of the thick

cm−2 were found. High-quality InAs

cm<sup>2</sup>

/V·s at 50 K.

The significant step in the development of HgCdTe photodiodes has been made by Arias with co-authors (Arias, 1993). They proposed the double-layer planar heterostructure (DLPH) photodiodes. The photodiodes were realized by incorporating a buried narrowbandgap active layer in the DLPH configuration. The planar devices were formed using a Hg1–yCdyTe/Hg1–xCdxTe (y>x) heterostructure grown by MBE on CdZnTe substrate. An im‐ portant feature of the DLPH approach is a planar p-doped/n-doped device geometry that in‐ cludes a wide-bandgap cap layer over a narrow-bandgap 'base' layer. The formation of planar photodiodes was achieved by selective implanting of arsenic through a ZnS mask fol‐ lowed by diffusing the arsenic (by annealing at high temperature) through the cap layer into the narrow gap base layer. After that the structures were annealed under Hg overpressure. The first high-temperature annealing was carried out to diffuse the arsenic into the base lay‐ er and to make the doped region p-type by substitution of arsenic atoms on the Te sub-lat‐ tice, while the second low-temperature one was carried out to annihilate Hg vacancies formed in the HgCdTe lattice during high-temperature process. In DLPH photodiodes sig‐ nificant reduction in tunneling current and surface generation-recombination current has been achieved. The architecture of mesa and planar heterostructure photodiodes are shown in Fig.1. The back-illuminated heterostructure photodiodes prepared from both LPE and MBE material grown on CdZnTe substrates have the highest performance achieved to-day (Arias, 1993; Bajaj, 2000; Rogalski, 2000).

The serious disadvantage of CdZnTe substrates is the thermal expansion coefficient mis‐ match with Si used for the read-out integrated circuit. To overcome this problem, growth of HgCdTe on alternate substrates such as sapphire, Si and GaAs has been developed. LPE, MBE and MOCVD techniques were used for this purpose. However, these alternate substrate ma‐ terials suffer from a large lattice mismatch with HgCdTe, leading to a higher defect density in the HgCdTe material and consequently reducing the detector performance (Bajaj, 2000). De‐ spite these disadvantages the large, 1024-1024 and 2048-2048, HgCdTe FPAs operating in SWIR and MWIR spectral bands were grown on alternative substrates (Kozlowski, 1999; Ba‐ jaj, 2000; Golding, 2003; Tribolet, 2003). Their performance is comparable to performance of FPAs prepared on lattice matched bulk substrates, with the same spectral cut-off.

Currently the major efforts are focused on photodiodes grown by MBE and MOCVD meth‐ ods. An overview *of* these technologies has been done by Razeghi (Razeghi, 2003). The InSb photodiodes were grown on 3-inch Si and (111)GaAs substrates. The InSb photodiodes typi‐

(n = 1015 cm−3 at 77 K) and a ~0.5 μm p+ (~1018 cm−3) contact layer. The crystallinity of these structure was excellent as confirmed by the X-ray diffraction, which showed FWHM <100 arcs for structures grown on (111)GaAs and Si substrates. Photodiodes were fabricated with 400 × 400 μm2 mesa structures by photolithography and wet chemical etching. These devices showed excellent response of about 1000 V/W, which is comparable to that of bulk detectors, with detectivities of ~3 1010 cmHz1/2/W at 77 K. It was shown that photodiodes can operate

A miniaturized InSb photovoltaic infrared sensor that operates at room temperature was developed by Kuze with co-authors (Kuze, 2007). The InSb sensor consists of an InSb p+

composition and thickness of AlInSb barrier layer was found to be x=0.17 and 20 nm, re‐ spectively. Typical responsivity of 1.9kV/W, output noise of 0.15 μV/Hz1/2 and detectivity of 2.8 108 cm Hz1/2/W was measured at 300K. InSb high-speed photodetectors were grown on semi-insulating GaAs substrate with 0.1 μm GaSb buffer layer using MBE (Kimukin, 2003). After the buffer layer a 1.5 μm thick n-InSb layer, 1.5 μm thick n-InSb layer, and finally 0.5 μm thick p*-* InSb layer were grown. The n-active layer was unintentionally doped to 2-3 1015 cm-3. Tellurium and beryllium were used as the n- and p-layer dopants, respectively. Doping level was 1018 cm-3 for both highly doped layers to decrease the serial resistance. The developed photodetectors can operate at room temperature. The responsivity 1.3 A/W was measured at 1.55 μm wavelength at room temperature. InSb photodiodes grown by MBE on GaAs coated Si substrates have been reported (Besikci, 2000; Ozer and Besikci, 2003). The peak detectivity of ~1·1010 cm Hz1/2 W−1 at 80 K has been measured under back‐ side illumination without anti-reflection coating. Differential resistance at 80 K is shown to be limited by ohmic leakage under small reverse bias and trap assisted tunneling under

InAs photodiodes are mainly manufactured by ion implantation and diffusion methods (McNally, 1970; Astahov, 1992). The p-n junctions prepared by Cd diffusion into n-InAs sin‐ gle crystals were investigated by Tetyorkin et al. (Tetyorkin, 2011). The photodiodes grown by MBE on InAs substrates have been reported (Kuan, 1996; Lin, 1997). The fabrication of InAs photodiodes on GaAs and GaAs-coated Si substrates by MBE have also been reported

In conclusion, during the past ten years the impressive progress has been made in the growth of narrow-gap semiconductors by MBE, MOVPE and MOCVD on silicon, GaAs and sapphire substrates. It is expected that in the near future HgCdTe and InSb IR photodi‐ odes will be grown directly on silicon for the low cost detectors with bi-spectral and multi-

up to room temperature, even though they were not optimized for this purpose.

structure grown on semi-insulating GaAs (100) substrate, with a p+

region (~1018 cm−3 at 77 K), a ~6 μm unintentionally doped region

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

layers to reduce diffusion of photoexcited electrons. The optimum Al



http://dx.doi.org/10.5772/52930

cally consisted of a 2 μm n+

er between p+ and p-

moderately large reverse bias.

(Dobbelaere, 1992).

spectral capability.

n+

**Figure 1.** Schematic cross section of mesa (left) and planar (right) HgCdTe IR photodiodes. The photodiodes are illumi‐ nated through the wide bandgap substrate.

While ion implantation of As requires activation of the implanted atoms at relatively high temperatures, an alternative technology (ion milling or plasma induced type conversion) has received considerable attention during the past few years. Currently plasma induced type conversion in HgCdTe is regarded as an alternative to ion implantation junction forma‐ tion technology (Agnihotri, 2002). The post-implant annealing is not needed in this technolo‐ gy. Reactive ion etching (RIE) induced type conversion and junction formation have been observed in a vacancy doped p-HgCdTe using H2/CH4 plasma. The junction depth could be adjusted from 2 to 20 μm.

#### *3.2.2. InSb and InAs photodiodes*

High performance InSb detectors have been fabricated with bulk material for decades. Typi‐ cally, the p-on-n junctions are prepared on bulk crystals of n-type conductivity with electron concentration in the range 1014 - 1015 cm-3 and mobility of the order of (2-6) 105 cm2 V-1 s-1 at T = 77 K. Be ion implantation and thermal diffusion of Zn and Cd seems to be the most fre‐ quently used technological methods which allow to obtain sharp p+ -n junctions (Mozzi and Lavine, 1970; Hurwitz and Donnelly, 1975; Rosbeck, 1981; Nishitani, 1983; Fujisada, 1985). The current status of InSb photodiode technology have presented by Wimmers et al. (Wimmers, 1983; Wimmers, 1988). Current manufacturing device processes require that bulk materials should be thinned before or after the junction preparation. Using this techni‐ que, hybrid FPAs were produced (Fowler, 1996; Hoffman, 1991; Parrish 1991). An array size of 1024-1024 was possible because the InSb detector material was thinned to <10 μm (after surface passivation and hybridization to a readout chip) which allows it to accommodate the InSb/silicon thermal mismatch. Certainly, the substrate thinning process is a very delicate process that can lower the yield and the reproducibility of the photodiodes, but it is used commercially because of LPE of InSb is not developed. Several attempts have been made to manufacture InSb photodiodes by LPE (Kosogov and Perevyaskin, 1970; Kazaki, 1976).

Currently the major efforts are focused on photodiodes grown by MBE and MOCVD meth‐ ods. An overview *of* these technologies has been done by Razeghi (Razeghi, 2003). The InSb photodiodes were grown on 3-inch Si and (111)GaAs substrates. The InSb photodiodes typi‐ cally consisted of a 2 μm n+ region (~1018 cm−3 at 77 K), a ~6 μm unintentionally doped region (n = 1015 cm−3 at 77 K) and a ~0.5 μm p+ (~1018 cm−3) contact layer. The crystallinity of these structure was excellent as confirmed by the X-ray diffraction, which showed FWHM <100 arcs for structures grown on (111)GaAs and Si substrates. Photodiodes were fabricated with 400 × 400 μm2 mesa structures by photolithography and wet chemical etching. These devices showed excellent response of about 1000 V/W, which is comparable to that of bulk detectors, with detectivities of ~3 1010 cmHz1/2/W at 77 K. It was shown that photodiodes can operate up to room temperature, even though they were not optimized for this purpose.

SWIR and MWIR spectral bands were grown on alternative substrates (Kozlowski, 1999; Ba‐ jaj, 2000; Golding, 2003; Tribolet, 2003). Their performance is comparable to performance of

**Figure 1.** Schematic cross section of mesa (left) and planar (right) HgCdTe IR photodiodes. The photodiodes are illumi‐

While ion implantation of As requires activation of the implanted atoms at relatively high temperatures, an alternative technology (ion milling or plasma induced type conversion) has received considerable attention during the past few years. Currently plasma induced type conversion in HgCdTe is regarded as an alternative to ion implantation junction forma‐ tion technology (Agnihotri, 2002). The post-implant annealing is not needed in this technolo‐ gy. Reactive ion etching (RIE) induced type conversion and junction formation have been observed in a vacancy doped p-HgCdTe using H2/CH4 plasma. The junction depth could be

High performance InSb detectors have been fabricated with bulk material for decades. Typi‐ cally, the p-on-n junctions are prepared on bulk crystals of n-type conductivity with electron

= 77 K. Be ion implantation and thermal diffusion of Zn and Cd seems to be the most fre‐

Lavine, 1970; Hurwitz and Donnelly, 1975; Rosbeck, 1981; Nishitani, 1983; Fujisada, 1985). The current status of InSb photodiode technology have presented by Wimmers et al. (Wimmers, 1983; Wimmers, 1988). Current manufacturing device processes require that bulk materials should be thinned before or after the junction preparation. Using this techni‐ que, hybrid FPAs were produced (Fowler, 1996; Hoffman, 1991; Parrish 1991). An array size of 1024-1024 was possible because the InSb detector material was thinned to <10 μm (after surface passivation and hybridization to a readout chip) which allows it to accommodate the InSb/silicon thermal mismatch. Certainly, the substrate thinning process is a very delicate process that can lower the yield and the reproducibility of the photodiodes, but it is used commercially because of LPE of InSb is not developed. Several attempts have been made to manufacture InSb photodiodes by LPE (Kosogov and Perevyaskin, 1970; Kazaki, 1976).

cm2


V-1 s-1 at T

concentration in the range 1014 - 1015 cm-3 and mobility of the order of (2-6) 105

quently used technological methods which allow to obtain sharp p+

nated through the wide bandgap substrate.

230 Photodiodes - From Fundamentals to Applications

adjusted from 2 to 20 μm.

*3.2.2. InSb and InAs photodiodes*

FPAs prepared on lattice matched bulk substrates, with the same spectral cut-off.

A miniaturized InSb photovoltaic infrared sensor that operates at room temperature was developed by Kuze with co-authors (Kuze, 2007). The InSb sensor consists of an InSb p+ -p- n+ structure grown on semi-insulating GaAs (100) substrate, with a p+ -AlxIn1-xSb barrier lay‐ er between p+ and players to reduce diffusion of photoexcited electrons. The optimum Al composition and thickness of AlInSb barrier layer was found to be x=0.17 and 20 nm, re‐ spectively. Typical responsivity of 1.9kV/W, output noise of 0.15 μV/Hz1/2 and detectivity of 2.8 108 cm Hz1/2/W was measured at 300K. InSb high-speed photodetectors were grown on semi-insulating GaAs substrate with 0.1 μm GaSb buffer layer using MBE (Kimukin, 2003). After the buffer layer a 1.5 μm thick n-InSb layer, 1.5 μm thick n-InSb layer, and finally 0.5 μm thick p*-* InSb layer were grown. The n-active layer was unintentionally doped to 2-3 1015 cm-3. Tellurium and beryllium were used as the n- and p-layer dopants, respectively. Doping level was 1018 cm-3 for both highly doped layers to decrease the serial resistance. The developed photodetectors can operate at room temperature. The responsivity 1.3 A/W was measured at 1.55 μm wavelength at room temperature. InSb photodiodes grown by MBE on GaAs coated Si substrates have been reported (Besikci, 2000; Ozer and Besikci, 2003). The peak detectivity of ~1·1010 cm Hz1/2 W−1 at 80 K has been measured under back‐ side illumination without anti-reflection coating. Differential resistance at 80 K is shown to be limited by ohmic leakage under small reverse bias and trap assisted tunneling under moderately large reverse bias.

InAs photodiodes are mainly manufactured by ion implantation and diffusion methods (McNally, 1970; Astahov, 1992). The p-n junctions prepared by Cd diffusion into n-InAs sin‐ gle crystals were investigated by Tetyorkin et al. (Tetyorkin, 2011). The photodiodes grown by MBE on InAs substrates have been reported (Kuan, 1996; Lin, 1997). The fabrication of InAs photodiodes on GaAs and GaAs-coated Si substrates by MBE have also been reported (Dobbelaere, 1992).

In conclusion, during the past ten years the impressive progress has been made in the growth of narrow-gap semiconductors by MBE, MOVPE and MOCVD on silicon, GaAs and sapphire substrates. It is expected that in the near future HgCdTe and InSb IR photodi‐ odes will be grown directly on silicon for the low cost detectors with bi-spectral and multispectral capability.

#### **3.3. Thermal annealing**

In HgCdTe photodiodes prepared on bulk crystals and LPE films the active absorption re‐ gion should be p-type with the hole concentration of the order of 1016 cm-3. Because of the asgrown materials are strongly p-type with vacancy concentration p>1017 cm-3, they should be annealed. Due to importance of annealing, it has been intensively studied in HgCdTe (Cap‐ per, 1994; Capper, 1997; Capper 2011). The most important parameters of annealing are tem‐ perature, Hg vapor pressure and annealing duration. In HgCdTe (x = 0*.*2) Vydyanath determined the free hole concentration at 77 K as a function of the Hg partial pressure, PHg, at anneal temperatures between 140 and 655 0 C. The concentration of doubly-ionized vacan‐ cies [VHg] is given by (Vydyanath, 1991):

At last, high temperature anneal in Hg vapor is used to activate the dopant by substituting

mediately after the high temperature anneal, annihilates the Hg vacancies formed in the

Bulk, LPE, MOCVD and MBE grown high purity material subjected to low-temperature

Techniques for improving the control of the Hg vacancy concentration was also reported by Yang et al. (1985), who deduced experimentally a relationship between the concentration of Hg vacancies, the annealing temperature, and the temperature of a Hg source. They also de‐ rived theoretically an analytic expression for this relationship. The annealing for MBE grown material can be conducted in-situ under ultra-vacuum conditions. In-situ annealing is very useful for production purposes. The improved ex-situ and in-situ annealing process‐

The performance of LWIR photodiodes can be improved by post-implant annealing (Bubu‐ lac, 1998). The dramatic decrease of the dark current in LWIR photodiodes was observed due to low temperature annealing at 120-150 °C (Ajisawa and Oda, 1995). The improve‐ ments was explained by changes in both carrier concentration profile and p-n junction posi‐ tion determined by interaction of interstitial Hg atoms with vacancies in the vicinity of the

Very scarce data are available in the literature concerning the annealing of InSb and InAs, both materials and devices. The effect of rapid thermal annealing and sulfur passivation on the quality of reactively sputtered SiO2 on InAs were investigated. Results show that both rapid thermal processing and sulfur passivation cause a reduction in leakage current

tion. Also, passivation started to loose its effectiveness when structures are annealed at

Tunneling current was observed by many authors in IR photodiodes made of narrow-gap A2B6 and A3B5 semiconductors (Rogalski, 1995). However, its nature seems to be under‐ stood in rare cases. For instance, the trap-assisted tunneling (TAT) via single level in the gap introduced by point defects was proved to be the main reason for the excess current in HgCdTe IR photodides at rather small reverse biases followed by the direct band-toband (BTB) tunneling current at higher biases (Nemirovsky, 1992; Rosenfeld, 1992; He, 1996). In these photodiodes the trap-assisted tunneling current is shown to be a source of

C) Hg reach annealing are converted to n-type with the electron concentration which depends on the doping level of residual donors. In HgCdTe material for LWIR photodiodes the minimum electron concentration that can be obtained is ~(2-10) 1014 cm-3 and the mobili‐

C), performed im‐

233

http://dx.doi.org/10.5772/52930

Infrared Photodiodes on II-VI and III-V Narrow-Gap Semiconductors

C caused degrada‐

arsenic atoms on the Te sublattice. A lower temperature anneal (200–250 0

HgCdTe lattice during high temperature treatment.

/V s (Chu and Sher, 2008).

es were developed for MOCVD grown films (Madejczyk, 2005).

and oxide fixed charges. Annealing at temperature higher than 400 0

**4. Carrier transport and recombination mechanisms**

**4.1. Tunnelling current in HgCdTe, InAs and InSb photodiodes**

(<300 0

ty is ~6 104

cm2

junction during the annealing process.

500 °C (Eftekhari, 1997).

[VHg]PHg= 1.7E28 exp(-1.67 eV/kT)cm-3atm

The hole concentration at 77 K is assumed to be equal to the concentration of vacancies. It was proved that there is no essential difference in the annealing behavior of bulk crystals and LPE epilayers. As seen, the same vacancy concentration can be obtained by combination of temperature and Hg vapor pressure. Obviously, samples annealed at lower temperature and pressure conditions would also have the lower concentration of interstitials and poten‐ tially the higher minority carrier lifetime. However, to reduce density of Te precipitates and dislocations high-temperature annealing should be performed (Capper, 2011). In epilayers grown on lattice-matched substrates (CdTe, CdZnTe) the dislocation density is in low-104 to mid-105 cm-2. At the same time, in epilayers grown on alternative substrates the dislocation density exceeds 106 cm-2. In LWIR epilayers grown by MBE on (211) GaAs and Si substrates the dislocation densities as low as 2.3 105 cm−2 have been obtained using high-temperature cycling between 300 and 490 °C (Shin, 1992). It has been shown that epilayers grown by MOCVD require a higher thermal annealing temperature than MBE material and the differ‐ ence in dislocation reduction between MBE and MOCVD HgCdTe materials is caused by dislocation movement under high-temperature and thermal stress conditions. A strong cor‐ relation between minority carrier lifetime and dislocation density was observed. The effect of temperature cycling on the dislocation density was also investigated by Farrell et al. (Far‐ rell, 2011). In-situ and ex-situ thermal cycle annealing methods have been used to decrease dislocation density in CdTe and HgCdTe. During the MBE growth of the CdTe buffer layer, the growth was interrupted and the layer was subjected to an annealing cycle within the growth chamber under tellurium overpressure. During the annealing cycle the temperature is raised to beyond the growth temperature (290 → 550 °C) and then allowed to cool before resuming growth again. This process was repeated several times during the growth. The insitu thermal cycle annealing resulted in almost a two order of magnitude reduction in the dislocation density. The decrease of dislocation density was attributed to the movement of the dislocations during the annealing cycles and their subsequent interaction and annihila‐ tion. To decrease the dislocation density in HgCdTe layers grown on CdTe/Si composite substrates, ex-situ annealing has been performed in a sealed quartz ampoule under a mercu‐ ry overpressure. It was found that the primary parameters that affect dislocation density re‐ duction are the annealing temperature and the number of annealing cycles.

At last, high temperature anneal in Hg vapor is used to activate the dopant by substituting arsenic atoms on the Te sublattice. A lower temperature anneal (200–250 0 C), performed im‐ mediately after the high temperature anneal, annihilates the Hg vacancies formed in the HgCdTe lattice during high temperature treatment.

**3.3. Thermal annealing**

232 Photodiodes - From Fundamentals to Applications

at anneal temperatures between 140 and 655 0

cies [VHg] is given by (Vydyanath, 1991):

[VHg]PHg= 1.7E28 exp(-1.67 eV/kT)cm-3atm

In HgCdTe photodiodes prepared on bulk crystals and LPE films the active absorption re‐ gion should be p-type with the hole concentration of the order of 1016 cm-3. Because of the asgrown materials are strongly p-type with vacancy concentration p>1017 cm-3, they should be annealed. Due to importance of annealing, it has been intensively studied in HgCdTe (Cap‐ per, 1994; Capper, 1997; Capper 2011). The most important parameters of annealing are tem‐ perature, Hg vapor pressure and annealing duration. In HgCdTe (x = 0*.*2) Vydyanath determined the free hole concentration at 77 K as a function of the Hg partial pressure, PHg,

The hole concentration at 77 K is assumed to be equal to the concentration of vacancies. It was proved that there is no essential difference in the annealing behavior of bulk crystals and LPE epilayers. As seen, the same vacancy concentration can be obtained by combination of temperature and Hg vapor pressure. Obviously, samples annealed at lower temperature and pressure conditions would also have the lower concentration of interstitials and poten‐ tially the higher minority carrier lifetime. However, to reduce density of Te precipitates and dislocations high-temperature annealing should be performed (Capper, 2011). In epilayers grown on lattice-matched substrates (CdTe, CdZnTe) the dislocation density is in low-104

mid-105 cm-2. At the same time, in epilayers grown on alternative substrates the dislocation density exceeds 106 cm-2. In LWIR epilayers grown by MBE on (211) GaAs and Si substrates the dislocation densities as low as 2.3 105 cm−2 have been obtained using high-temperature cycling between 300 and 490 °C (Shin, 1992). It has been shown that epilayers grown by MOCVD require a higher thermal annealing temperature than MBE material and the differ‐ ence in dislocation reduction between MBE and MOCVD HgCdTe materials is caused by dislocation movement under high-temperature and thermal stress conditions. A strong cor‐ relation between minority carrier lifetime and dislocation density was observed. The effect of temperature cycling on the dislocation density was also investigated by Farrell et al. (Far‐ rell, 2011). In-situ and ex-situ thermal cycle annealing methods have been used to decrease dislocation density in CdTe and HgCdTe. During the MBE growth of the CdTe buffer layer, the growth was interrupted and the layer was subjected to an annealing cycle within the growth chamber under tellurium overpressure. During the annealing cycle the temperature is raised to beyond the growth temperature (290 → 550 °C) and then allowed to cool before resuming growth again. This process was repeated several times during the growth. The insitu thermal cycle annealing resulted in almost a two order of magnitude reduction in the dislocation density. The decrease of dislocation density was attributed to the movement of the dislocations during the annealing cycles and their subsequent interaction and annihila‐ tion. To decrease the dislocation density in HgCdTe layers grown on CdTe/Si composite substrates, ex-situ annealing has been performed in a sealed quartz ampoule under a mercu‐ ry overpressure. It was found that the primary parameters that affect dislocation density re‐

duction are the annealing temperature and the number of annealing cycles.

C. The concentration of doubly-ionized vacan‐

to

Bulk, LPE, MOCVD and MBE grown high purity material subjected to low-temperature (<300 0 C) Hg reach annealing are converted to n-type with the electron concentration which depends on the doping level of residual donors. In HgCdTe material for LWIR photodiodes the minimum electron concentration that can be obtained is ~(2-10) 1014 cm-3 and the mobili‐ ty is ~6 104 cm2 /V s (Chu and Sher, 2008).

Techniques for improving the control of the Hg vacancy concentration was also reported by Yang et al. (1985), who deduced experimentally a relationship between the concentration of Hg vacancies, the annealing temperature, and the temperature of a Hg source. They also de‐ rived theoretically an analytic expression for this relationship. The annealing for MBE grown material can be conducted in-situ under ultra-vacuum conditions. In-situ annealing is very useful for production purposes. The improved ex-situ and in-situ annealing process‐ es were developed for MOCVD grown films (Madejczyk, 2005).

The performance of LWIR photodiodes can be improved by post-implant annealing (Bubu‐ lac, 1998). The dramatic decrease of the dark current in LWIR photodiodes was observed due to low temperature annealing at 120-150 °C (Ajisawa and Oda, 1995). The improve‐ ments was explained by changes in both carrier concentration profile and p-n junction posi‐ tion determined by interaction of interstitial Hg atoms with vacancies in the vicinity of the junction during the annealing process.

Very scarce data are available in the literature concerning the annealing of InSb and InAs, both materials and devices. The effect of rapid thermal annealing and sulfur passivation on the quality of reactively sputtered SiO2 on InAs were investigated. Results show that both rapid thermal processing and sulfur passivation cause a reduction in leakage current and oxide fixed charges. Annealing at temperature higher than 400 0 C caused degrada‐ tion. Also, passivation started to loose its effectiveness when structures are annealed at 500 °C (Eftekhari, 1997).
