Metallurgy of ohmic contacts to InP

Metallurgy of ohmic contacts to InP

Canadian MeraNurgrrd Quarter/~, Vol. 34, No. 2, pp. 85-l 13, IY95 Copyright I’; 1995 Canadian Institute of Mmng and Metallurgy Printed in Great Bntam...

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Canadian MeraNurgrrd Quarter/~, Vol. 34, No. 2, pp. 85-l 13, IY95 Copyright I’; 1995 Canadian Institute of Mmng and Metallurgy Printed in Great Bntam. All rights reserved 000X4433/95 $9 SO+O.OO

Pergamon 000&-4433(94)00026-3




DOUGLAS Department

of Mining,


Metallurgical Edmonton,





and Petroleum Engineering, Alberta, T6G 2G6, Canada


of Alberta,

Abstract-Metallurgical aspects in the design and formation of ohmic contacts to InP are considered and reviewed in this paper. InP and related ternary and quaternary semiconductors (e.g. InGaAs and InGaAsP) are utilized primarily in optoelectronic devices found in fibre optic applications, such as lasers and photodetectors. These devices are operated under high current densities (4-10 kA/cm*), with dimensions in the range 2-150 pm in width. Low resistance (ohmic) contacts are required to link the active regions of the semiconductor devices to the external circuit. The small dimensions associated with these devices place severe materials processing constraints on the formation of ohmic contacts. Some of the issues that have to be addressed include contact stability (the metallization/semiconductor system is a multicomponent system in a nonequilibrium state), metallization adhesion to the semiconductor, contact uniformity and extent of reaction with the semiconductor. A very powerful tool for examining contact metallizations in three dimensions on a very fine scale is transmission electron microscopy (TEM). This technique can provide imaging, structural and compositional information from very small volumes. The capabilities of TEM analysis in characterizing metallizationjsemiconductor microstructures are demonstrated, with emphasis on the above issues.


cet article, nous r&visions les aspects m&allurgiques dans la conception et la formation de

contacts ohmiques avec I’InP. Les InP et autres semi-conducteurs tertiaires et quaternaires (InGaAs et InGaAsP) sent surtout utilisCs dans des dispositifs optoClectroniques que I’on retrouve dans des applications de fibres optiques comme les lasers et les photodttecteurs. Ces dispositifs fonctionnent sous des courants de haute densite (de 4-10 kA/cm*) avec des dimensions allant de 2-5 pm en largeur. Des contacts (ohmiques) de faible resistance sont n&cessaires pour relier les rtgions actives du dispositif semi-conducteur aux circuits externes. Les petites dimensions allant de pair avec ces dispositifs sont g l’origine de restrictions dans le traitement de la formation des contacts ohmiques. Nous devons considtrer certains problkmes comme la stabilitC des contacts (le systime de mtttallisation/semi-conducteurs un systtme & plusieurs composantes hors d’equilibre), l’adhttsion de la mCtallisation au semi-conducteur, l’uniformiti: des contacts et Y&endue des r&actions avec le semi-conducteur. La microscopic B transmission ilectronique (TEM) est un outil trts puissant pour ttudier la mttallisation des contacts en trois dimensions sur une tchelle tr.?s rtduite. Cette technique peut fournir des informations imagbes, structurelles et compositionnelles g partir de tr&s petits volumes. On a surtout dCmontrt. dans cet article, les capacitts analytiques d’un TEM pour caracttriser la microstructure m&tallisation/les semi-

conducteurs, en insistant sur les questions &oqu&es ci-dessus. INTRODUCTION

high contact resistance with respect to the electronic devices; an ohmic contact has a linear current-voltage characteristic due to a small contact resistance relative to the device resistance. Ohmic contacts on InP-based materials are necessary for such electronic devices as field effect transistors (FET) [ 1,2], junction field effect transistors (JFET) [3,4], high electron mobility transistors (HEMT) and heterojunction bipolar transistors (HBT) [5, 61, as well as for photonic devices like long-wavelength laser diodes, light-emitting diodes (LED) [7-91 and photoelectronic and solar cells [lo]. Semiconductor devices are being continually scaled down in size and operated at higher current densities and elevated temperatures, placing ever-increasing demands on ohmic contact performance. The main requirements for ohmic contacts are low contact resistance and good thermal stability, which are directly affected by the contact microstructure, as well as good adhesion, uniformity and a shallow reaction zone at the metal/semiconductor interface. It is essential to understand the relationship between the microstructure and the electrical properties of the metal/

InP, a direct band gap compound semiconductor, has become an important semiconductor material because of its unique combination of electrical and optical properties, e.g. high electron mobility, high saturation velocity and radiation tolerance. The primary application of InP and related alloys in the semiconductor industry is in the fabrication of optoelectronic devices or lightwave devices. During device operation, electrical communication links between the active regions of semiconductor devices and the external circuit are provided by metal/semiconductor contacts. Generally, when a metal and a semiconductor are brought together, there exists an energy barrier for electron transport from the semiconductor side to the metal side at the metal/ semiconductor interface, due to the difference in work function between the metal and semiconductor, According to their current-voltage characteristics, contacts can be divided into two categories. A Schottky barrier or rectifying contact has a 85



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semiconductor contacts in order to fabricate high quality and reliable ohmic contacts. A great deal of work has been done and reported in the literature (e.g. Refs [1 l]-[l4]) regarding the formation and performance of ohmic contacts to GaAs, another III-V semiconductor. Somewhat less has been done, until recently, on contacts to InP. In particular, detailed analysis of the microstructural changes that occur during ohmic contact processing and how these changes affect contact resistance, stability and reliability has been lacking. A major reason for this has to do with problems associated with microstructural characterization. Transmission electron microscopy (TEM) is one of the best techniques available to monitor and analyze microstructural changes on the scale required for ohmic contacts. Through the use of both plan-view and cross-section specimens, three-dimensional structural and compositional information is attainable with a spatial resolution not available with other analytical techniques such as scanning electron microscopy @EM), Auger electron spectroscopy (AES), secondary ion mass spectroscopy (SIMS) and x-ray diffraction (XRD). One of the drawbacks of TEM lies in sample preparation, particularly the preparation of cross sections of metallizations on InP. This is probably the major reason for a lack of detailed microstructural analysis of contacts to InP. It is the aim of this paper to examine ohmic contacts to InP from a metallurgist’s point of view. With this in mind, the paper will be divided into several parts, starting with a brief description of some of the properties and applications of InP. This will be followed by a discussion on metal/semiconductor contacts, including the formation and characterization of ohmic contacts. The bulk of the paper will deal with actual contact metallizations (both single and multilayer) to InP, with an attempt to define the role(s) of each constituent (if there is one) in the formation of an ohmic contact.










Heavy 7



holes (hh


Light holes (lh)


Split-off holes (soh) Fig. l(a). E-K diagram for a direct bandgap semiconductor showing three hole masses: heavy holes, light holes and split-off holes [16].


Semiconductors have electrical resistivity values that lie between conductors and insulators, i.e. from about 10m2 to IO9 Qcm [15]. There are two basic types of semiconductors-the more common elemental semiconductors (e.g. Si and Ge) and compound semiconductors made up of two or more elements. One of the most important groups of compound semiconductor is the IIIIVs, which consist of at least one Group III element (such as Al, Ga and In) and at least one Group V element (such as P, As and Sb). The atoms in III-V semiconductors are tetrahedrally bonded, commonly in a zinc-blended type structure. GaAs and InP are two common examples of III-V semiconductors. The energy band structure of solid materials consists of a valence band and a conduction band. For most semiconductors, the valence band structures are similar, consisting of a heavyhole (hh) and a light-hole (lh) band, because of their large and small effective masses, respectively, and a split-off hole (soh) band as shown in Fig. 1 [ 161. The maximum in the valence band is located at K = 0 (K represents momentum space), i.e. the center of the Brillouin Zone. The conduction band structures for the various semiconductors are different. The conduction band structure is comprised of several minima, including a



holes (soh)

Fig. l(b). E-K diagram for an indirect bandgap semiconductor showing that the minimum in the conduction band occurs at K away from K = 0 [16].




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single minimum at K = 0 and additional minima along the [ 1001 and [l 1 l] axes. The band gap is the difference in energy between the lowest point of the conduction band and the highest point of the valence band. If the conduction band minimum at K = 0 is lowest in energy as shown in Fig. l(a), the semiconductor is a direct band gap material ; otherwise, it is an indirect band gap semiconductor, as shown in Fig. l(b). InP, GaAs and related ternaries and quaternaries (e.g. InGaAs and InGaAsP) are direct band gap semiconductors. The mobility and effective mass of electrons in a semiconductor strongly depend on the band gap structure of the semiconductor. In a direct band gap semiconductor, an electron in the conduction band can fail directly to an empty state in the valence band, giving off the energy difference E, as a photon of light. On the other hand, an electron in the conduction band minimum of an indirect semiconductor cannot fall directly to the valence band maximum but must go through a state within the band gap. In an indirect transition, which involves a change in K, the energy is lost largely as heat to the lattice rather than as an emitted photon. The mobility of electrons (p) is higher and the effective mass of electrons (m*) is lower in direct band gap semiconductors than in indirect band gap semiconductors. The difference between direct and indirect band structures is very important in determining which semiconductors can be used in devices requiring light output. The calculated energy band structure of InP is shown in Fig. 2 [17]. Both the valence maximum and the lowest conduction band minimum occur at the center of the Brillouin Zone (f point). The conduction band minimum is doubly degenerate and is separated by an energy E, from the valence band maximum. Away from the zone center, the highest valence bands separate into two sets of bands, referred to as the heavyand light-hole bands. The next set of valence bands, referred to as the spin-split-off bands, also have their maxima at I, where they are doubly degenerate separated by an energy A from the maxima of hh and Ih bands. Higher conduction band minima




occur in the [ 1001 and [ 11 l] directions, as can been seen in Fig. 2. The minima in the [ll I] and [loo] direction are, respectively, 0.4 eV and 0.7 eV above the I minimum. InP has become an important semiconductor material because of its unique combination of electrical and optical properties. The properties of InP are compared with those of more common semiconductor materials, GaAs and Si, in Table 1 [16, 18-201. For high-speed devices with electrons as the majority carriers, a semiconductor with high electron mobility and high saturation velocity is desirable. High thermal conductivity and high breakdown fields are also desirable properties. All of these requirements are well met in InP [21, 221. Its many alloy forms, such as InGaAs or InGaAsP, provide a variety of energy gaps and carrier mobilities. The primary use of InP and its alloys with Ga and As in the semiconductor industry is in the fabrication of optoelectronic devices or lightwave devices for fiber optic communication, such as lasers, photodetectors, solar cells. photodiodes and optoelectronic integrated circuits. In addition, much attention has been given to InP due to its tolerance to radiation. Compared with Si and GaAs, InP is superior in a radiation environment. At present, InP is still not as popular as Si or GaAs in industrial applications due to its fragility, thermal stability and sensitivity to ion-induced damage compared with Si and GaAs [23]. The binary nature of the InP lattice makes ion implantation damage removal and dopant activation steps more complex than for Si. Amorphization of the In-based semiconductor during implantation and incongruent evaporation of P from the sample surface upon high-temperature annealing are also serious problems for InP. In addition, because InP and its alloys are more expensive than Si and GaAs, device production costs are several hundred times as high.

METAL/SEMICONDUCTORCONTACTS Based on their current-voltage characteristics, metal/semiconductor contacts can be divided into two classifications. Contacts with rectifying characteristics are called Schottky barriers or rectifying contacts, and contacts with linear characteristics are referred to as ohmic contacts or nonrectifying contacts.


o’4 eV $-*eV


E,= 1.34eV

Schottky contacts


Figure 3 [24] shows an energy-band diagram of a metal and an n-type semiconductor separated by a large distance. In the figure, I$,,, indicates the work function of the metal, i.e. the energy needed to remove an electron from the Fermi level of the metal (ET) to the vacuum level. The work function of the semiconductor is & and 4. = xs+s’>

Fig. 2. Calculated


band structure


of InP in two principal

tions in the Brillouin Zone [ 171.



where the electron affinity xs is the difference in energy between an electron at the vacuum level and an electron at the bottom of the conduction band and r is the difference in energy between the Fermi level in the semiconductor (E‘r) and the bottom of the conduction band. In the case where the semiconductor is n-type and 4, < 4, (Fig. 3) if the metal and the semiconductor are brought together



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1. Properties



of InP, GaAs



and Si [IS, 17-191


(300 K)

Energygap (F,) (ev) Electrons

Effective mass (m*) Holes (hh) Holes (hl) Intrinsic mobility (cm’/V-s) Electrons Holes Resistivity p (Rem) Electron affinity x (v) Crystal structure Lattice constant (nm) Density (g/cm’) Thermal expansion coefficient Thermal conductivity (W/cm) Melting point (“C)



1.34 0.073

1.42 0.007

1.12 m, = 0.97 m, = 0.19

0.4 0.08

0.45 0.08

0.49 0.16

8500 400 4X 10% 4.07 Zinc blend 0.565 5.32 5.9 x 1omh 0.46 1238

Vd = hn-A

Valence Band

Fig. 3. Energy-band diagrams of a metal and an n-type semiconductor separated by a large distance. The work function for the metal is &,, (see Ref. [24]).

(the vacuum level is assumed to be the same in both materials), electrons flow from the semiconductor to the metal, accumulating at the surface of the metal and depleting the number of electrons in the surface region of the semiconductor. The resulting dipole electric field (due to the positive ions) opposes further electron flow, and in equilibrium,


In the Schottky model, the semiconductor is assumed to be homogeneous right up to the boundary with the metal, so that the uncompensated donors give rise to a uniform space charge region of depth w, called the depletion depth or depletion region [Fig. 4(a)] [24, 251. The induced electric field strength E, therefore, increases linearly with distance from the edge of the depletion region [Fig. 4(b)], and the electrostatic potential 9 increases quadratically [Fig. 4(c)]. The resulting parabolic barrier is known as a Schottky barrier, as shown in Fig. 4(d). The depletion layer width MI is an inverse measure of the square root of the doping concentration [ 181. The energy barrier c#J~~is the energy difference between the Fermi level in the metal and the bottom of the conduction band in the semiconductor at the interface. 4en = 4m-XS>


1500 450 2.5 x 10’ 4.05 Diamond 0.543 2.33 2.6 x lo-” 145 1420

where &,n is the Schottky barrier height of an n-type semiconductor, i.e. the energy necessary for electrons in the metal to penetrate to the semiconductor, and it is the single most important parameter describing a Schottky barrier. The barrier for electrons in the semiconductor relative to the metal is V,, the diffusion potential or the band bending in the semiconductor at equilibrium :


EF = EL.


4600 150 2X 10’ 4.38 Zinc blend 0.587 4.79 4.56 x 10m6 0.68 1070

Conduction Band




= dh-SI.


The current-voltage characteristics of a metal/semiconductor contact are governed by the transport of the charge carriers (i.e. electrons or holes) across the metal/semiconductor interface and its associated space-charge region. For a Schottky barrier, the current transport mechanism is thermionic emission. The majority carriers, with kinetic energies in excess of Vd. are emitted from the semiconductor over the top of the barrier into the metal. If a negative voltage V is applied to the semiconductor, the semiconductor/metal barrier is reduced to Vd - V while & remains unchanged. The reduction of the barrier on the semiconductor side makes it easier for electrons in the semiconductor to move to the metal. This is the forwardbias condition, and a large current can flow. On the other hand, if a positive voltage is applied to the semiconductor, the potential barrier is increased to prevent current conduction and reverse bias is achieved. The current-voltage relationship for an n-type semiconductor is an exponentially increasing current under forward bias and almost constant current under reverse bias. This rectifying characteristic is shown in Fig. 5 [16]. The energy band diagram of a metal and a p-type semiconductor pair, with &,, < & before contact is shown in Fig. 6(a). When the metal and semiconductor are joined, the positive carriers from the semiconductor cross over to the metal until the Fermi level of the metal-semiconductor system is aligned. This results in the band diagram depicted in Fig. 6(b). Similar to the metal/n-type semiconductor case described before, there are a Schottky barrier height &,, for current (or positive carriers) from metal to semiconductor and a barrier Vd for current from the semiconductor to metal. When a forward bias V is applied to the contact (p-type semiconductor positive with


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V ReWP Fig. 5. Schottky barrier I-V characteristics for a circular deposited on an n-type semiconductor [16].




the crystal lattice at the interface produces a large number of energy states, called interface or surface states, located within the forbidden gap. The interface states are usually continuously distributed in energy and are characterized by a neutral level E,,. The states below E, are positively charged when empty ; the states above EO are negatively charged when occupied. If E,, is

Vacuum Level

0s---_.f. 1 -/_E IT

ConductJon Band

Cd) Fig. 4. Variation of (a) charge density, (b) electric-field strength, and (c) electrostatic potential with distance from the metal/semiconductor interface according to the depletion approximation. (d) Band diagram of an ideal metal/n-type semiconductor interface [24. 251.

Valence Band



(a) to metal), which raises the electrostatic potential on the semiconductor side relative to the metal side, the barrier V, is lowered to a smaller value ( Vd - v). Current can more easily flow from semiconductor to metal. For a reverse bias (- I’), the opposite occurs, the electrostatic potential of the semiconductor side is depressed relative to the metal side, and the potential barrier becomes larger (V, + I’). Therefore, it is difficult for positive carriers to pass the barrier from semiconductor to metal, i.e. there is no current flow in the contact. The contact thus shows a rectifying characteristic. For a given semiconductor and for any metal, the sum of the barrier height on n-type and p-type substrates is equal to the energy band gap [26], or: respect

Fig. 6(a). Energy-band

diagrams of a metal and ap-type separated by a large distance.

Vacuum ,/--



Level -

(b) Fig. 6(b). Band diagram

In practical metal/semiconductor

contacts, the disruption


of an ideal p-type face.




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in the depletion layer to reach equilibrium. As a result, the diffusion potential V, is effectively reduced, and according to equation (4) the barrier height &n is also reduced [Fig. 7(a)] [18]. Similarly, if E,, < E,, there is a negative charge in the interface states and & is increased to bring EF closer to E, again. Thus, the change in the interface states has a negative feedback effect which tends to keep EF closer to E. [Fig. 7(b)]. If the interface state density is large, the Fermi level is effectively pinned at E, and &rn becomes independent of the work function of the metal and semiconductor. In most practical Schottky barriers. interface states play a dominant role in determining the value of &,,, and the barrier height is essentially independent of both the work function difference and doping level in the semiconductor. Ohmic

0) Fig. 7. Pinning of the Fermi level by surface states. All energy states below EF are occupied [5]. (a) E0 is above E,, (b) E0 is below E,.

aligned with the Fermi level, the net charge of the surface is zero. Whenever E. > E,, the net charge of the interface states is positive, or donorlike, so that fewer ionized donors are needed


If a metal and an n-type semiconductor pair with #Jo < c#J,, the exchange of carriers leading to a constant Fermi level results in the band diagram depicted in Fig. 8(b) [24]. The potential barrier at the junction is almost nonexistent, so that carriers can freely pass in either direction. As a result, this metal/ semiconductor contact is ohmic. Also, it is shown in Fig. 8(c) that a metal/p-type semiconductor with &, > & is an ohmic contact. In addition to the zero potential barrier case, there are other possibilities for forming ohmic contacts in metal/semiconductor junctions. The width of the depletion layer is inversely dependent on the square root of the doping level, so that the probability of quantum-mechanical tunneling of charge carriers through the barrier increases as the doping level increases. At very high doping levels, i.e. with impurity densities (NJ of lOI cm-’ or more, the barrier can be thin enough to permit electrons (or holes), with energies close to the Fermi energy in the semiconductor, to tunnel through the bar-

a. Rectiying

b. Ohmic

c. ohmic

d. Rectifying

Fig. 8. Band diagrams illustrating the Schottky theory of the metal/semiconductor contact are possible depending on the difference 4,,-& and the semiconductor majority-carrier current flow [24].

interface. Four type. The arrows

types of indicate


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TlX%mionic-field emission Field emission

Fig. 9. Field and thermionic-field emission under forward bias. The diagram refers to a degeneratively doped n-type semiconductor for which 5 is negative [25].

rier at low temperature. This is known as “field emission”. At moderately high doping levels (10” cm-3 < Nd < lOI cmm3 [27]), the barrier is somewhat wider. However, if the temperature is increased, electrons (or holes) can still gain enough thermal energy (E,) and tunnel near the top of the barrier producing an appreciable current. This process is known as thermionic-field emission. Figure 9 [25] shows both field and thermionic-field emission for a metal on an n-type semiconductor. Field emission is independent of temperature, while thermionic-field emission is temperature dependent. Both types of quantum-mechanical tunneling can lead to ohmic behavior in the contacts. The current-voltage relationship of ohmic contacts is approximately linear as shown in Fig. 10 [ 181. When ohmic contacts are applied to uniform semiconductor samples or devices, the measured currenttvoltage relationship is determined by the resistivity of the semiconductor sample or by the behavior of the device to which the contact is formed, rather than by the characteristics of the contact. Practically speaking it is not essential that the current-voltage characteristics of the contact itself be linear, provided that its resistance is very small compared with the resistance of the specimen or device. In addition: the contact should not inject minority carriers and should be stable both electrically and mechanically. Formation

qf ohmic contacts

Modern high-speed electronic and photonic devices are operated under high current densities (410 kA,/cm’). with dimensions in the range 22150 pm in width [23]. They require a very low specific contact resistance for the contact which links the active region of the semiconductor device to the external circuit, in order that only a negligible voltage drop occurs during operation. From the preceeding discussion, several possibilities exist for forming low resistance ohmic contacts : (1) A metal can be chosen such that its work function is less than the work function of an n-type semiconductor or greater than the work function of ap-type semiconductor. This is usually not possible for InP. (2) A semiconductor with an extremely narrow band gap. such as InAs (E, = 0.35 eV), which can be epitaxially grown and lattice matched to the semiconductor is another possibility. InAs has high electron and hole mobilities (33,000 and 460 cm’,V.s, respectively, at 300 K) and has surface states pinned in the conduction band ; therefore, an ohmic contact can be formed with virtually any metal that is deposited onto it [23].


Fig. 10. I-V characteristics of an ohmic contact [18].

(3) A third possibility is to provide a very heavily doped (> 5 x 10”cmm3) thin layer on the semiconductor surface, which is to be immediately adjacent to the metal. The depletion region is then so thin that field emission may take place and the contact has a very low resistance at zero bias. This can be done in two ways. One method involves growing a heavily doped epitaxial layer on the semiconductor prior to metal deposition. Often, the epitaxial layer is a lower bandgap material such as In, ,,Ga, 47A~ (E, = 0.75 eV), which is lattice matched to InP or a graded In,Ga, -,As layer (graded to InAs) [23]. These contacts should not, at least in theory, require subsequent annealing and are therefore referred to as non-alloyed contacts. Practically speaking, most of these contacts are still annealed to attain optimum contact resistance values. The other way of producing doped surfaces is through an external dopant source. The dopant is usually driven into the semiconductor by means of diffusion, which requires heating of the contact. These contacts are usually termed alloyed contacts. Doping could also be achieved by ion implantation of the appropriate species; however, this leads to problems such as surface damage and loss of stoichiometry [23]. Of the above methods

listed above for fabricating





contacts, the most common for III-V semiconductors is the third method, i.e. introducing a highly doped semiconductor region just below the metallization. Our discussion on the metallurgical behaviour of ohmic contacts will then almost exclusively focus on this process. In designing contacts to InP (or other III-Vs), a number of factors have to be considered. The most important of these is probably the contact resistance, although several other important aspects have to considered, and these have been summarized by Katz [23] : (1) The contact should have the lowest possible resistance. (2) The contact should be stable over a wide temperature range. Limited interfacial reactions and contact uniformity are desirable. (3) The contact should be stable throughout device electrical and thermal operating conditions. (4) Stresses in the metal films should be minimized. (5) The metal(s) used in the contact structure should be compatible with the metals used in the interconnects. (6) The fabrication process should fit as an integrated step into the overall device manufacturing scheme. It is, of course, difficult to meet all these requirements in a given contact metallization. One of the most difficult constraints to fulfill is maintaining contact stability and uniformity while limiting interfacial reactions. This will be a major consideration in the following discussion. Characterization

of ohmic contacts

Contact resistance measurement. The most commonly used parameter to electrically characterize an ohmic contact is the specific contact resistance rcr defined by

(6) where V is the voltage and J is the current density. For a homogeneous contact of area A with uniform current density, the contact resistance R, is simply Rc=A.



tacts to n-type compared to p-type InP, due to higher barrier heights (0.7-0.8 eV for metals on p-type InP vs 0.40.5 eV for n-type InP) and larger carrier effective masses for the p-type material [23, 241. Microstructural characterization. Several techniques can and have been used to characterize contact microstructures. Some of these, as pointed out in the introduction, include SEM, TEM, XRD, AES and SIMS. All of these provide useful information, but they have their limitations. AES and SIMS, for example, provide a means of doing depth profiling on the nanometer scale, but have poor lateral spatial resolution. Of the available techniques, we find that the best way to obtain a three-dimensional picture is through TEM analysis of cross-section and plan-view specimens. TEM is not without its limitations, however. The sample volumes analyzed are generally quite small (of the order of 1 pm3), although the volumes do exceed the dimensions of the microstructures to be investigated. The biggest limitation is probably sample preparation, particularly cross section samples. This will be discussed in some detail in the following paragraphs. TEM characterization. Two types of samples are generally prepared for TEM analysis, i.e. cross-section and plan-view specimens. Cross-section specimens are particularly useful as they allow imaging parallel to the metal/semiconductor interface. Plan-view specimens permit imaging perpendicular to the interface and generally provide a larger field of view. Plan view specimens are relatively easy to prepare [Fig. 1 l(a)]. A 3-mm disc is cut from the wafer, followed by mechanical polishing from the semiconductor side down to a thickness of about 50 pm. Final thinning to perforation is achieved by single-sided chemical jet polishing using a solution of 5% Br in ethanol. In many cases the metallizations consist of two or more metal layers, before and/or after annealing. It may be beneficial to isolate one of the layers from the others to facilitate phase analysis. This is achieved by sputtering at low energy (334 keV)

rc (7)

The measured resistance R will be approximately equal to R, for most sample geometries when r, > lo-’ Rem’. However, for smaller values of rcr the spreading resistance of the semiconductor R, and the series resistance R, of the connecting wires and semiconductor substrate must be taken into account. Then. in general, R = R,+R,+R,,


where R, and R, depend on the particular geometry of the metal/semiconductor contact being characterized. For III-V semiconductors, accurate determination of r, is most often carried out using one of four methods [24] : (1) The Cox and Strack technique; (2) the four point method; (3) Shockley extrapolation technique; or (4) the transmission-line model. These techniques are discussed in detail in Ref. [24] and will not be discussed further here. Specific contact resistance values as low as = 1O-’ Rem* can be measured with an error of about k 25%. Contact resistance values are generally lower for ohmic con-

(b) glue’

cross section specimen

” semiconductor

Fig. 11. Schematics depicting preparation of (a) plan-view and (b) crosssection specimens.


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from either the semiconductor or the metallization side for a brief period of time (of the order of a few minutes). By using several samples and a combination of chemical polishing and sputtering, it is possible to prepare thin sections of each layer of interest. Cross-section specimens are somewhat more difficult to prepare. The procedure we use [28] is a variation of the technique originally described by Bravman and Sinclair [29]. Pieces approximately 5 mm long and 1 mm wide are cleaved from the metallized wafer. Two pieces are glued together face-to-face and then glued to pieces of scrap InP wafer to form a “raftlike” structure [Fig. 1l(b)]. The pieces are glued such that the two sample pieces are in the middle of the raft. A 3-mm diameter disc is cut from the raft and then polished with 600 grit SIC paper to a thickness of 200 pm. One side is polished with 1 ,um diamond paste to a mirror finish. The other side is dimpled to near perforation (central thickness is about 20 pm) and then polished with 1 pm diamond paste to a mirror finish and to produce a small hole less than 0.5 mm in diameter. The final stage is to ion mill or sputter the specimen, for 3045 min, at low energy (34 keV). The specimen is cooled during ion milling in order to minimize preferential sputtering effects. InP tends to sputter much more rapidly than the metal layers, making it difficult to produce electron transparent regions of all areas of interest. In addition, InP itself is subject to preferential sputtering problems-P tends to sputter faster than the In, leaving globs or islands of In on the surface [30]. Both problems are minimized by liquid nitrogen cooling. Dimpling the samples to perforation also helps to minimize the preferential sputtering problem, as the time for ion milling is greatly reduced. However, it should be pointed out, that it can be difficult to dimple the samples to perforation without breaking the thin material around the hole. TEM is best done in an analytical TEM, also referred to as an analytical electron microscope or AEM, equipped with an energy dispersive x-ray (EDX) detector and capable of doing convergent beam electron diffraction (CBED). Both of these techniques provide high spatial resolution (typically better than 10 nm), with EDX analysis providing compositional information and CBED providing structural information. Further details on these techniques are beyond the scope of this paper, but they can be found in, for example, Ref. [3 11.Where possible, other analytical techniques, such as XRD, SEM, etc., are useful for corroboration purposes. In fact, we do XRD analysis on all of our samples prior to TEM to determine which samples to examine in the TEM. XRD indicates when phase changes have occurred.




As described in the section on the formation of ohmic contacts, the most common method of producing ohmic contacts is to highly dope the surface region of the semiconductor. In theory, this should be achievable by depositing a single layer of the appropriate metal. In practice, however, contacts are multilayer structures, consisting of, for example, a dopant metal, a diffusion barrier and a nonoxidizing capping layer. The result, from a metallurgical point of view, is an unstable,





multicomponent system, with a minimum of 3. and possibly as many as 6 or 7, elements. Reactions occur during annealing. leading to the formation of several intermetallic compounds between the various metals and the semiconductor constituents. Because of the complexities of the reactions between metals and InP, and the lack of necessary information such as phase diagrams for ternary and quaternary systems, thermodynamic data (heat of formation, free energy change of formation, etc.), kinetic data (diffusivities. reaction rates) and crystal structures. at present, the fabrication of contacts is still more of an art than a science. Most publications in this field report the experimental results for particular contact systems, i.e. the formation techniques, contact properties and, in some cases, limited microstructural information. In the following sections, we will examine and analyze published results on metal/InP systems. The discussion will focus primarily on Au-based and near-noble metal-based contacts to InP, emphasizing both the advantages and disadvantages of each. One particular system, i.e. Pd-Ge to n-type InP, will be examined in detail, with consideration given to optimization of the contact structure, from both an electrical and a metallurgical standpoint. Au-based metallizations Au-based metallizations are the oldest and most widely used ohmic contacts to III-V semiconductors. Au-based ohmic contacts have several beneficial features, i.e. relatively low specific contact resistance and low sheet resistance of the metallic layer, high resistance against corrosion, compatibility with standard evaporation and photolithography techniques and fairly good reliability at room temperature. Au-ZnP Contacts. Gold is used extensively as a metallization for compound semiconductors, such as GaAs [13, 32-361, GaP [3740], InSb [41], GaSb [41, 421 and InGaAsP [43, 441. In the Au-semiconductor contact systems mentioned above, Au acts as an oxidizing agent and displaces the group V elements in combining with group III elements to form various kinds of compounds. Similarly, in Au/InP contacts the dominant reaction in the ternary Au-In-P system is between Au and In, leading to formation of Au-In compounds. There are several binary thermodynamically stable Au-In phases, but only one thermodynamically stable Au-P phase (Au?P,) [47] (summarized in Table 2 [45-48]). The reactions between Au thin films and InP single crystal substrates have been reported by several groups [37, 38,49961]. The metallurgical reactions occurring in the Au/InP system are presented schematically in Fig. 12(a). The thermodynamic nature and the highly reactive nature of InP make reactions between Au and InP possible at very low temperature. Several groups [42,49951] have reported that during gold deposition at room temperature, Au reacts with both In and P. In and P out-diffuse substantially into the metallization with different diffusion rates. Hiraki [49] reported, using AES, a fairly large quantity of In atoms on the surface of a Au film (100 nm thick) just after deposition, and P atoms at the surface 100 min after deposition. During annealing at 250&45O”C, even though different AuIn compounds have been reported by various research groups due to different processing conditions, the reaction process can




Table 2. Au-In intermetallic compounds [45-48] Phase

at% In


me (“C)

AH (kJ/mole Au)





0 to -8.78

; I 61

12-14.3 13-23 21.5-22.2

492-641 641-649

- 8.78 13.54toto - -9.20 17.56 - 16.72


21.7-22.5 24.5-25 29.8-30.6 28.8-31.4

481492 457487

- 16.72 to - 17.97 -21.74

e AuIn

35.3-39.5 50-50.1

Au8h3 Au,In Au,In2 Au,,In, Au&, Au,In Au,ln, Au,ln, Au,In Au& AuIn

454458 510

-28.84 to -31.35 -44.73





- 76.08

be roughly divided into three stages : (1) Au-In solid solution formation, (2) Au,In and Au2P, formation and (3) Au,InAu&, transition. Each stage is controlled by a different mechanism. Initially (< 32O”C), interaction between Au and InP begins with the entry of In into Au via a dissociative diffusion process. Dissociative diffusion is characterized by the rapid diffusion of In via an interstitial mechanism through the Au layer until substitution occurs on vacant Au lattice sites, which are generated primarily at the surface. This can result in a maximum in In concentration profiles near the contact surface. In continues to dissolve into Au until the solid solubility limit is reached. This process is controlled by the vacancy generation rate at the free surface of the metallization. The P atoms, released when In initially enters the metal, remain unreacted and either leave the system (through sublimation) or occupy nonlattice sites in the metal. The first step proceeds very rapidly even at room temperature, but can be affected by the surface condition of the contact metallization. For instance, a SiO, capping layer on the free surface of the gold film suppresses the metal/semiconductor reaction rate by inhibiting the vacancy generation rate at that surface [57]. In the second stage (32&37O”C), individual islands of Au,In (or Au+,) nucleate and grow more or less uniformly over the

as &posited

Pearson symbol

Structure type

Lattice parameter @ml



a = 0.4106

hP16 hP2 Hexagonal OC28 hP26 OP8 hP60 CP52

a a a a a a a

hP5 Triclinic

Ni,Ti Mg Cu,,Sb Cud% /?-Cu,Ti Au& Cu,Al, N&Al3 -



= = = = = = =

a= a= c( = a=

0.2909, c = 0.95 0.2917, 0.4798 1.053, c = 0.4788 0.4784, b = 2.0437, c = 0.5071 1.0539, c = 0.4786 0.5864, b = 0.4745, c = 0.5165 1.2215, c = 0.8509 0.9829 0.4538, c = 0.5645 0.43, b = 1.059, c = 0.356 90.54”. /? = 90’, y = 90.17” 0.6507

metallization layer. The highly mobile In diffuses interstitially from the InP/metal interface to the Au&/Au-In solid solution interface, and occupies substitutional sites in the metal lattice by displacing Au atoms. This is the so called “kick-out” mechanism that results in an elevated concentration of highly mobile interstitial Au. The displaced Au atoms become substitutional again by annihilating the vacancies that have been generated at the InP/metal interface during In diffusion into the metal. The Au atoms then react with unbound P, remaining there to form interfacial Au*P,. The rate-limiting step in this stage is the release of In from the InP and its insertion interstitially into the Au lattice. Indium entry continues until the metallization is converted from the saturated Au(In) solid solution to Au&. In the final stage, Au31n transforms to Au&, (or Au&). This transition is controlled by an Au-In exchange or kick-out process which takes place at the interface between the Au,In phase and the Au&, phase. At the same time, the P released during this stage leaves the system unreacted. In a closed Au/InP system, for a thin layer of Au [system composition is designated as x in Fig. 12(b)] [51], three phases coexist at equilibrium, i.e. Au& (or AuJn), Au,P, and InP. But for an open Au/InP system (which is often the case for practical contacts), gas phase P will be liberated from the

320- 370Oc

370 - 4oo*c

Fig. 12(a). Schematic representation of reactions between Au films and InP substrates.

400 - 45oOc


and PING



Fig. 12(b). Au-In-P

phase diagram [51]

InP type

200 nm Au


300 nm(Au-Ge) 100 nm(Au-Ge)/ 100 nm Au

n n n n

250 nm(Au-Sn) nm Ge/80 nm Au/50 nm Tij300 nm Au/40 nm Zn/250 nm Au nm Zn/200 nm Au nm Au/40 nm h/270 nm Au

urn Au

250 nm (Au-Be) 80 nm (Au-Be)/210

nm Au

100 nm(Au-Be) 100 nm(Au-Be)/200

nm Au

20 nm Mn/ 150 nm Au 40 nm Mgi 160 nm Au 40 nm Ni/ 140 nm Au

n P P P P P P P P P

60 nm AuBe/SO nm Cr/300 nm Au nm Cr/200

nm Au

25 nm AuZn/SO nm TiWjlOOO nm Au



Au-based/InP contucts. Au-based multilayer metallization schemes usually consist of an external dopant (either p- or ntype) source, which is included with the intention of alloying with the adjacent semiconductor through solid or liquid phase reaction during annealing to form a p+ or n+ region. This highly doped region, according to the simple model discussed above, should then reduce the width of the depletion layer associated with the potential barrier formed at the metal/semiconductor interface, making the contact essentially ohmic. Au, in these types of contact, is either deposited as the top layer of the contact or as a binary mixture with the dopant species. Gold provides a non-oxidizing protective layer during annealing, and appears to play a role in lowering the contact resistance. Values of specific contact resistance (rJ for several Au-based contacts are summarized in Table 3, and it is clear that in all cases contact resistances are significantly lower compared to those for pure Au contacts (comparisons are made for n-type InP only). This indicates that the additional elements have some effect, be it doping or otherwise, which will be discussed further

Table 3. Au-based contacts to InP Doping level Resistance


80 nm AuBe/20


A limited amount of work has been done on Ag metallizations [82-871. The annealing behaviour of these contacts is quite similar to Au metallizations and therefore a brief discussion is included in this section. The major difference is that the diffusion process is more unbalanced [82], i.e. In and P diffuse considerably faster into Ag than Ag diffuses into the semiconductor. The result is the formation of AgP, at the metallization/semiconductor interface and a Ag-In solid solution. Unlike the Au/InP system, no In binary compounds form, at least in the temperature range studied (30&45O”C). Ohmic contacts, with resistances in the low lO-6 Rem’ range, have been formed with 200 nm thick Ag layers on n-type InP substrates, with initial doping levels of 1.7 x 10’” crne3 [82]. This value is about an order of magnitude better than that achieved for Au/InP contacts with similar doping levels. The onset of ohmic behaviour has been attributed to the formation of AgP, [82].

system, with the evaporation of P increasing abruptly above 550°C [62, 631. The local composition near the contact area becomes more In-rich. Indium rich Au-In compounds (AuIn? or Au&,) are then observed, with no additional Au,P, forming in the system. As deposited, Au contacts on n-type InP are Schottky barriers with a barrier height of 0.4-0.5 eV [64]. The Schottky barrier remains stable on annealing up to 350°C. Above this temperature, the Schottky barrier is catastrophically degraded due to Au-In and Au2P, compound formation. A large 2-3 order of magnitude drop in the contact resistance occurs when the Au/InP contact is annealed at 400°C [37]. At this point, the character of the contact changes from Schottky to ohmic [65] (see e.g. Table 3).

45 40 65 20



(cm-‘) 1 x 10’8 3 x 1o’8 1.7 x0’: 1.6 x 10” 3 x 1o’8 1 x loI* 5.5 x 18’” 10’” 5.5 x 10’8 5.5 x lO1” 1 x lOI8 1.5 x 10” 1.3 x 10’8 5.5 x 10” 2x lOI 2 x lOI 6 x 10” 2 x 10’” 5.5 x 10” 2 x lOI I x 1o’B 1 x lOI

(km’) 5 x 10-S 1 x lo-’ 0.07 Rmm 0.03 Qmm 1.8 x 10m6 4 x 10-l 2x lo-’ 9.4 x lo-> 1.8 x lO-4 2.2 x 1om4 <8x 1o-5 1 x 1om3 2 x 1om4 7x1o-4 4.5 x 1o-4 6 x 1O-4 1 x 10-d 9 x 1om4 1 x 1om4 3.5 x lo-4 5 x 1om5 3.7 x 1om5


Processing condition 4OO”C, 60 s 375x, 30 s 42O”C, 3&270 s

Reference 38 66 67

30 s 90 s 3 min 6 min 3 min 3 min 10 min

68 69 70 71 72 72 73 74

42&43O”C 3-4 min 45oq I5 s 446°C. 50 min 4OO”C, 30 s 42O”C, 4 min 42O”C, 3 min 41O”C, 34min 410°C. 30 s


39o”c, 48o”C, 42O’C, 43O”C, 42O”C, 46O”C, 42O”C,

76 77 78 79 80 81



and PING


in later sections. Au has a high solubility for In and the dopant element and can react with In and P to form several kinds of compounds, leading to InP decomposition, which is believed to be responsible in part for dopant incorporation. Au-based contacts require annealing to between 350°C and 48O’C in order to achieve ohmic behaviour. For contacts to p-type InP, Au-Zn [70-72, 88, 901 and AuBe [72-75, 79,90,91] metallizations have been commonly used. In the Au-Zn/InP system, the metallization layers are made by either depositing Au and Zn metal layers separately (indicated here as Au/Zn, where Zn is adjacent to the semiconductor) or co-depositing a Au and Zn mixture (indicated here as AuZn). Zn reacts with Au to form AuZn during deposition. On annealing up to 340°C both Au and Zn diffuse towards the substrate and In diffuses outward to the metallization. It has been found that Au diffuses into InP more rapidly than Zn does and Au reacts with both Zn and P. AuZn transforms to Au,Zn before Au-In phase formation (about 36OC). The contact resistance decreases and the contact becomes ohmic at about 380°C. The optimum specific resistance (r, = 1.8 x 10-&m* with an initial InP doping level of 5.5 x 10” cmp3 [72]) is associated with the formation of a Au-In phase (Au,In, or Au&,) and Au2PI in the temperature range 42C~48O”C. It is postulated that Zn occupies the lattice sites created by In out-diffusion, to create a layer of higher carrier concentration required for achieving lower contact resistance, and also segregates in the Au2P, clusters. At higher temperatures (about SIOC), ZnP, forms and contact resistance increases again. Because of the large difference in the vapor pressure of Au and Zn, and a very low sticking coefficient of Zn on III-V semiconductor surfaces, specific technological problems associated with the deposition and alloying of Au-Zn contacts can arise. Typical examples of the problems are preferential evaporation of Zn prior to Au deposition, poor adhesion, nonhomogeneity of Au-Zn depositions and re-evaporation of Zn during the alloying of Au-Zn metallizations. As a result, AuBe metallizations have been studied as alternatives. The interactions between Au, Be, In and P in the Au-Be/InP system are very similar to those in the Au-Zn/InP system except no AuBe and Be-P compounds form. To obtain ohmic properties and low specific contact resistance values, it is necessary to alloy the contact at temperatures above 400°C. At about 42045O”C, the contact is composed of Au&, and Au,P, phases. Contact resistances as low as 8 x lo-’ Rem’ have been reported for ptype InP substrates with an initial doping level of 1 x 10’8cm~3 [73]. The Au to Be ratio appears to influence the optimum heat treatment time and temperature, but not the specific contact resistance. The minimum value of Y, is independent of the method of sample preparation and the amount of dopant alloyed into the InP surface with the Au film. In addition to Au-Zn and Au-Be metallizations, Au-Mn [76], Au-Mg [77] and Au-Ni [78] systems have been applied to p-type InP. The specific contact resistances for these contacts are 6 x 10m4 Qcm’ with an initial doping level of 2 x 10’” cmm3, 1 x lop4 Rcm2 with an initial doping level of 4 x lOI cm-’ and 9 x 10m4 Qcm’ with an initial doping level of 2 x 10” cm-), respectively. Au-Ge metallizations are commonly used for n-type InP [66, 67,90, 921. The metallizations are based on a Au-12.5 wt% Ge eutectic alloy. The Au-Ge contact was originally chosen for its




low eutectic melting temperature (361°C). Alloying it at the eutectic melting temperature was then thought to cause localized melting of the Au-Ge and drive the dissolution of InP. Indium would then diffuse outward into the Au and Ge would be incorporated preferentially on In sites, as a donor. When Au-Ge/InP is annealed at temperatures between 320°C and 450°C however, all of the reactions are solid state in nature. This is not surprising since the metallization is no longer a AuGe binary layer, but consists of several binary phases (e.g. GeP and AuJn). The phase change sequence in the metallization is : GeP + GeP + Au,In + GeP + Au,In + Au,P, + GeP + Au&, + Au2P,. The contact becomes ohmic after annealing at 320°C. At 375-4OO’C the lowest specific contact resistance (Y, = 1 x lo-’ 0cm2, with an initial doping level of 3 x lOI cmm3) is observed [66]. Annealing the contact at 450°C and higher results in significant out-diffusion of P and enhanced formation of P vacancies. This leads to degradation of the contact electrical properties. It should be emphasized that no melting occurs during annealing, i.e. all reactions are solid state in nature. Au-Sn is another metallization that is used for n-type InP contacts [68, 79, 931. Au-Sn/InP contacts exhibit ohmic behaviour in a narrow alloying temperature range between 380°C and 410°C. The minimum specific contact resistance, rc = 1.8 x 10d6 Qcm’, is achieved by annealing at 390°C for 30 s [68]. Au,In, AuSn and polycrystalline InP are the main phases in the contact at alloying temperatures between 380°C and 410°C. Above 410°C the contact is predominantly composed of AuJn. It seems that the reaction processes for Au-based/InP contacts are similar to that of Au/InP contacts, except for the formation of additional binary phases, containing the dopant species, i.e. M-Au, M-P or M-In (M = Zn, Sn or Ge). Au-In and Au,P, binary phases form in the same temperature range as for the Au/InP system. The lowest values of specific contact resistance for Au-based/InP contacts are obtained in those samples annealed in the temperature range 39042O”C, which is also the optimum temperature range for Au/InP contact processing. Comparison of values of specific contact resistance (Table 3) for n-type InP specimens with the same doping level, shows that the r, values of Au-based contacts are l-2 orders of magnitude lower than those for pure Au contacts. This implies that the additional elements, i.e. Sn or Ge, are important in Au-based contacts. The major problem associated with Au-based contacts is gold’s affinity for the Group III element (In in InP). Gold diffuses to the InP surface and drives the out-diffusion of In. Gold diffusion to the semiconductor surface leads to the formation of Au-In compounds (as discussed above), which tend to grow unevenly or “spike” into the InP substrate. Examples of this phenomenon will be shown for multilayer metallizations in later sections. The outward diffusion of In to the Au layer compromises the long-term stability and integrity of the contact structure. Generally, a relatively thick Au layer (several hundred nm) is deposited on top of the Au-based metallizations, providing a low resistance contact to the rest of the circuit and protection from oxidation. The presence of Au, however, drives the undesired out-diffusion of P and the dopant element, which can degrade the electrical properties and decrease device reliability. In addition, there is evidence that contact resistance degrades over time, as much as 20% over a period of several weeks [80].



In order to alleviate this problem, another metal layer is often deposited between the top Au layer and the Au-based contact to serve as a diffusion barrier for P and the dopant element. For example, Au/Cr/Au-Be [SO, 911, Au/TiW/Au-Zn [81] and Au/Ti/AuGe[69] metallization systems have been studied. It has been shown that the diffusion barrier layer (Cr, TiW or Ti) not only prevents the out-diffusion of P from the semiconductor but also prevents the migration of Au from the top layer into the underlying semiconductor during heat treatment. With the diffusion barrier, the dopant remains at the contactjsemiconductor interface. The contacts do not degrade with time and device reliability is superior to those without diffusion barriers. The electrical properties of these contacts are also given in Table 3. Another remedy to this problem is to deposit the outer Au layer after contact annealing. Near-noble transition metal-based metallizations Near-noble transition metal-based metallizations are introduced into ohmic contact schemes in an effort to overcome the deficiencies involved in the processing and performance of Aubased contacts to both p-type and n-type semiconductors, particularly their unstable nature and lateral nonuniformity. Because of the more stable thermodynamic properties and less reactive nature of these metals, the near-noble transition metalbased metallizations provide a less reactive and thermally more stable contact microstructure during heat treatment at higher temperatures. Ohmic behaviour may be achieved through a combination of doping and the formation of narrow band gap intermetallic phases at the metal/semiconductor interface. Near-noble transition metal contacts. The first step toward understanding the behaviour of a complex near-noble metalbased metallization system is to investigate the interaction between single metal layer and compound semiconductor substrates and to determine the role of the near-noble metal in the metallization system. The commonly used near noble transition metals to III-Vs are Ni and Pd, and to a lesser extent Co, Rh, Pt and Ir. Common features of these metals include ease of deposition, low reaction temperatures with III-V compound semiconductors (which greatly enhances contact adhesion to the semiconductor), and resistance to oxidation and compound formation with both the Group III and Group V elements. Pd/InP and Ni/InP are examined in detail here, with some recent work on Pt/InP metallizations also discussed. The reactions that occur during annealing are quite similar for all three metallizations. These can be summarized as follows and are discussed below. (1) Pd, Ni and Pt react with InP at low temperatures to form amorphous ternary phases that are metal rich. This phenomenon is believed to be due to rapid diffusion of the metal into the semiconductor and subsequent disruption of semiconductor bonds. The composition of the amorphous phase lies within the composition range defined by a line drawn between the metal and InP on the phase diagram. (2) Annealing at intermediate temperatures results in crystallization of the amorphous phase, forming one or more ternary crystalline phases.






Fig. 13. Pd-In-P phase diagram at 600°C [94]. The Pd-rich portion of the diagram is omitted.

(3) Annealing at higher temperatures results in decomposition of the ternary phases into binary phases or elemental species. Portions of isothermal sections of the phase diagrams, determined from bulk specimens, for both the Pd-In-P [94] and NiIn-P [95,96] systems are shown in Figs 13 and 14. Dashed lines on the phase diagrams represent calculated tie lines. Note the similarities between the systems, which are reflected in the similarities in reaction sequences. In addition, all known binary and ternary compounds (including metastable phases) for these systems are listed in Tables 4 and 5. Pd/ZnP Metallizations. The PdjInP system is examined initially, and is used to demonstrate the capabilities and utility of TEM analysis in investigating the complex reactions that can occur in thin film metallizations. In a recent comprehensive study [103, 1091 on the reactions between Pd thin films (60 nm thick) and InP (annealed in vacua) three crystalline ternary phases, Pd*InP, Pd,InP and Pd21nP(II), as well as an amorphous ternary phase (Pd,,InP), were identified. (An amorphous phase of similar composition and Pd,InP were also reported to form in another study by Caron-Popowich et al. [ 1 lo].) The amorphous phase begins to form during Pd deposition and continues to grow during annealing up to 225°C. At this point, crystalline islands of cubic (Ll,-type structure) PdzInP begin to nucleate and grow at the amorphous layer/InP interface. This is clearly shown in the cross section and plan view micrographs in Fig. 15(a). The ternary phase grows epitaxially on InP according to the following epitaxial relationship : (lOO)Pd,InP]](llO)InP


Pd,InP grows into a continuous layer (4&50 nm thick) after annealing for 15 min at 225°C [Fig. 15(b)]. Pd,InP forms at this point at the Pd,InP particle boundaries and grows epitaxially with respect to Pd,InP :


and PING






to a third crystalline ternary phase, referred to in Ref. [103] as Pd,InP(II) since it has the same composition as PdJnP, but a different structure. The structure of Pd,InP(II) had not been determined previously but was characterized by means of electron diffraction analysis. It was found to have a similar crystal structure to PdJnP-the structure is cubic with double the lattice parameter of Pd,InP (a = 0.830 nm). The following epitaxial relationship with InP was reported in Ref. [109] : (l00)Pd,InP(II)/~(ll0)InP

Annealing of the Pd,InP samples at temperatures of 400°C or higher results in the decomposition of Pd,InP(II) into two binary phases, i.e. PdIn and PdP, (Fig. 17). These are the equilibrium phases in contact with InP, as long term annealing (350 h at 500°C) results in no further phase transformations. These results agree quite well with those expected from the bulk PdIn-P phase diagram (Fig. 13). The overall composition of the thin film couple is indicated by an X on the phase diagram, and lies within the three phase field bounded by InP, PdIn and PdP,. Stremsdoerfer et ul. [ 111, 1121 investigated reactions between Pd and n-InP in specimens where Pd films were deposited by electroless deposition, instead of the more conventional electron beam evaporation or sputtering techniques, and then annealed at temperatures up to 500°C. An optimum contact resistance of 5 x 10e7 &rn* (the initial doping level was 10” cm-‘) was attained after annealing at 300-350°C. XRD analysis was used to characterize the reaction products. InPd was identified in as-deposited samples. Pd,In and PdIn, were found at 300°C and only Pd,In was detected at 400°C. No phosphide or ternary phases were reported to form.


(b) 600°c





Fig. 14. Ni-In-P phase diagram at (a) 470°C and (b) 600°C. Tie lines not determined by experiment are denoted by dashed lines [95,96]. The Ni-rich portion of the diagram is omitted.



[1 lO]Pd,InP]][OOl]Pd,InP





It appears that PdJnP forms first because of kinetic considerations, i.e. it is closer in composition than Pd,InP to the amorphous phase (Pd,,InP). As Pd,InP grows, Pd is rejected and accumulates ahead of the PdJnP growing front, since Pd,InP has a lower Pd content relative to the amorphous phase. The last portions of the amorphous phase to crystallize are Pdrich, which results in the formation of PdJnP in the boundary regions (Fig. 16). The entire Pd layer is consumed by 275”C, and the amorphous layer has fully transformed to Pd,InP and Pd,InP (10-l 5% by volume). Annealing at 300-350°C results in the transformation

NijZnP metalhations. Nickel is a common metal in ohmic contacts to InP, not only because it is widely used in Au/Ge/Ni contacts to n-type InP, but also because it shows potential on its own as an ohmic contact to n-type InP. For instance, Appelbaum et al. [ 1131 have reported that evaporated N&P films exhibited low specific contact resistance on n-InP (rc = 3 x 10e6 Slcm’ for an original doping level of 2 x IO’* cme3). Several groups have investigated reactions involving Ni films on n-InP during annealing at 20&500°C in forming gas (5%H2, 95%N,) [111, 114, 1151, helium [113] and vacuum [108]. Although different annealing ambients were used, the results were quite similar. The results of the TEM study done by ourselves [IO81 are summarized below and can be compared with results from Ref. [114] (Fig. 18). As with the Pd/InP system, some of the Ni reacts with InP during deposition to form an amorphous ternary phase (7-9 nm thick). The formation of an amorphous ternary phase has been reported previously by Sands et al. [I 141 for low-temperature annealing (~200°C) of 40 nm Ni films on InP, and is the first reaction product in AuiNi contacts to InP [78]. Annealing the films at low temperatures results in growth of the amorphous layer, with complete consumption of the Ni by 250°C. The composition, as determined by EDX analysis, corresponds approximately to Ni,InP, which is very close to the value reported by Sands et al. (Ni, ,InP) [114]. The first signs of crystallization of the amorphous layer are apparent in samples annealed in the 25s300°C range. Several phases have been identified, including a ternary phase which appears to match the one reported by Sands et al. (NiJnP) [114] and two binary


and PING


Table 4. Pd-In-P





Structure type





a = 0.71067,







75.0-80.0 69.2

1047 814 860 1150

mP28 mC16 OP16 hR20 mC12 C132 t12 OP12 OP16 CP2 hP5 Cubic tP7 Tetragonal

a c a a a a

= = = = = =

0.5004, b = 0.7606, 0.8416, ,0 = 95.63” 0.5674, b = 0.94409, c = 0.821, fi = 110.41’ 0.2837, b = 0.9441, c = 0.7695, /? = 90.2” 0.5947, b = 0.7451, c = 0.517 1.1976, 6 = 0.7055

a u a a a a a a a a

= = = = = = = = = =

0.6207, b = 0.5857, c = 0.5874, 0.770s 0.407, c = 0.38 0.5611, b = 0.4218, c = 0.8237 0.56, b = 1.102, c = 0.424 0.326 0.453, c = 0.55 0.94144 0.393, c = 0.692 0.8376, c = 0.8158

at% Pd

DTI- (“C)


89.2 81.9

66.2-67.7 74.0-75.2 63.5-61.2 48-61.5 39-39.5 25.0

1304-1365 1066 1285 709 664

50.0 50.0

the decomposition

at% Ni

N&In Ni,In N&In &Ni-In Ni,& NiIn NiIn N&In Ni,In Ni,,In,, N&In3 Ni?81n77 Ni,In, NiIn, N&P Ni,P, Nir2P5 Ni,P N&P, NiP NiP, NiPZ NIP, N&InP


5. NikInP

Ni,Sn Au&, N&In

a a a a

= = = =

0.532, c = 0.424 0.375 0.4179, c = 0.5131 0.418, c = 0.512-0.518

a a c u a a

= = = = = =

0.4545, 0.3060 0.2929 0.74, b 1.4646, 0.4396,



CF116 Hexagonal

a a a a a a a a a a a a a a a

= = = = = = = = = = = = = = =

0.918 0.6199. 0.8954. 1.322, 0.8646, 0.5864. 0.6789, 0.605. 0.54706 0.6366, 0.80151 0.681, 0.681, 1.1112 0.412,

670 876 860

870 410

970 1175 1106 850

Amorphous the decomposition

47, 101, 102 47,101 47, 101 47,99 47, 101 99 97,99 47, 100 98 47,98 47 47,98 47,98 47,98 47 94 94 103 103, 109 103


hP8 CP4 hP6 Monoclinic hP4 hP6 CP2 CI2 mP22 mC44 hP5 Cl52 C140 t*12 t132 hP168 t134 hP9 hP36 OP16 CP12 mC12 C132 Monoclinic





66.7 62.0-66.97 59.5-62.5 49.5-50.5

Ni&P, NiJnd’,, Ni, ,InP Niz,-,InP

fl = 111.8’

a = 0.415 a = 0.830 -

Structure type


75.0 71.7 70.5


c = 1.70867

Pearson symbol



40.0-41.0 29.5-30.0

PdPTp, As&o In Co$i Ge,RhS ClCs Al,Ni, IX-3 Pd,TlAs

parameter (nm)




CFe, Pd,P,

Cubic Cubic Amorphous


PDT represents



Pearson symbol


Pd,P Pd,P Pd, iP Pd,P Pd,P Pd,P W’, PM’, PdPz PdP, Pd,In Pd,In PdJn, PdIn Pddn, PdIn, Pd,InP Pd,InzPl Pd, ,InP PdZInP Pd,InP(II) Pd,,InP



NiAs CoSn ClCs w In3Pt, A13NiZ ;a-Cu,Zn, GeJr, N&P Ni12P( Fe-P NijP, NiP Fe& PdP, A&o


parameter @ml

Reference 47, 104 47, 104 106 104 47, 104

c = 0.4353

= 0.426, c = 1.046, /i’ = 90” b = 0.8329, c = 0.8977, 0 = 35.35” c = 0.521

L’ = 0.6128 c = 0.4386 c = 2.4632 c = 0.507 c = 0.3385 c = I .0986 b = 0.4881, c = 0.689 b = 0.5615,

c = 0.6072,

fl = 126.22’

b = 0.529, c = 1.280, fl = 95’ b = 0.529, c = 0.640, fi = 94” c = 0.483

47 47 47 47, 104 91,104 47 47 47,105 47, 105 47, 10s 47, 10s 47. 105 47 47 47 95 108 97,107 95 96, 108



and PING







Fig. 15. Bright field TEM micrographs from cross-section and plan-view specimens from Pd/InP annealed at 225’C for (a) 5 min and (b) 15 min. Selected area diffraction (SAD) patterns are also shown along with the plan-view micrographs. The diffuse ring in the SAD patterns is from the amorphous phase while spots are from the Pd,InP in (a) and Pd,InP and Pd&P in (b) [103].

phases, NizP and In,O,. The crystallization process is. unlike the Pd/lnP system, very much laterally non-uniform. Two regions from the same specimen are shown in Fig. 19-one area is made up of Ni,lnP and a mixed region of Ni,InP and Ni2P, while the other area consists of the amorphous phase and In203. Ni,InP forms adjacent to the InP substrate and exhibits an orientation relationship with respect to the InP, i.e. [OOl] Ni,lnP 11[OOI] InP (indicating that NiJnP most likely nucleated initially at the amorphous layer/InP interface) [IO& 1141. The ternary

phase was identified as being monoclinic with a = 0.681 nm, h = 0.529 nm, c = 0.640 nm and b = 95”. The micrographs in Fig. 19 illustrate the heterogeneous nature of the reactions and also demonstrate the importance of examining both cross-section and plan-view specimens. Crosssection specimens permit observation of samples parallel to the interface, but only provide a very limited field of view. Planview specimens, on the other hand, provide a much better picture of the overall microstructure and phase distribution.



Fig. 16. Dark field image of a plan view specimen from Pd/InP annealed at 225°C for 15 min. Pd,InP particles show up bright. The dark field image was taken using a (002) reflection from Pd,InP [103].

Fig. 17. Bright field micrograph (cross section) from Pd/InP sample annealed at 500°C for 5 min [103]. Annealing at 5OOC results in partial decomposition of the ternary phase (NiJnP), resulting in additional Ni,P formation and release of metallic In :

Ni,InP --t Ni,P + In. <200*c




Some of the released In diffuses to the surface where it oxidizes to form In,O, and a complex In-P oxide, which indicates some loss of P as well. The main feature of the Ni,/InP reactions, and the major difference compared with the Pd/InP system, is the laterally non-uniform morphology of the crystalline phases. In particular, the ternary phase does not form continuously across the InP surface, indicating that Ni,InP formation is likely nucleationcontrolled. Similar behaviour has been reported by Appelbaum et al. [113], who have done Auger analysis of Ni/InP reactions. At 250°C (3&60 min), Auger results indicated an intermixed region of Ni, In and P. The intermixed region was likely the same as our amorphous phase (and the phase reported by Sands et al. [114]-annealed in forming gas). Appelbaum et al. [113] also report “phase separation” for annealing at times longer than 60 min at 250°C or at 300°C. Auger results indicate a NiP-rich region and a Ni-In-rich mixture with some P in the outer layer. Because of the nature of Auger analysis, Appelbaum et al. were unable to conclusively identify the phases. It is likely that the Ni-P-rich region is the same as our NizP and the Nip In-rich region may have been NiJnP. Sands et al. [ 1141 have reported similar results for Ni/InP couples annealed in forming gas. The major difference is the formation of a hexagonal ternary phase (Ni, ,InP), which formed prior to NiJnP (monoclinic) formation in their work. Fatemi and Weizer [115] report that two phosphides form during annealing at 400°C i.e. Ni,P and N&P. Longer-term annealing (-40 min) results in the decomposition of N&P, leaving only N&P. These phases were identified by EDX and AES analysis. Examination of the Ni-In-P phase diagram [Figs 14(a) and 14(b)] indicates that below 526’C, at the overall composition corresponding to the thin film couples (indicated by X in the phase diagram), InP and NiJnP should be the equilibrium phases. However, loss of In, due to oxidation (In,O,), would make the couple somewhat P-rich, leading to the formation of Ni,P (or possibly other more P-rich) phases. This is consistent with the thin film results described above. Above 526’C, the ternary phase is no longer stable, i.e. InP, NizP and In are the thermodynamically stable phases. The electrical properties of Ni/InP contacts consisting of only ternary phase morphology have not been reported, but for NiiInP contacts consisting of nickel phosphides. ohmic behaviour has been achieved. Appelbaum et al. [ 1161 investigated the interaction of sputtered N&P films on n-InP and p-InGaAs substrates and found that sputtered N&P layers form ohmic

200 - 3oo*c


Fig. 18. Schematic representation of the reactions of Ni films on InP substrates [114]

360 5OO*C


and PING





Fig 19. BI right field micrographs and diffraction patterns from a plan-view specimen of a sample an]nealed at i !80”( :f br S min. The micrograph in (a) shows a Ni,InP region and a mixed Ni,InP,/Ni,P regi Ion. while the micrograph in (b) shows an area with the amorphous phase and In20, [108]. contacts with low specific contact resistance on n-InP (3 x 1Om6 Rcm2, n = 2x 10” cm-‘) and p-InGaAs (2 x 10m5 Rem’, p = 8 x lOi cm-3). As a result, the Ni contact does not require additional dopants (Ge or Zn for n- or p-type, respectively) to become ohmic, thereby eliminating the need for a hightemperature electrical activation treatment. It has been proposed that the Ni/InP contact becomes ohmic because Ni2P forms a low barrier height contact to the underlying InP. It

should also be pointed out that the initial doping levels for these contacts is quite high, which may be a contributing factor. In addition, N&P should be stable on InP, at least at temperatures below 6OO”C, as there is a tie line joining the two phases on the phase diagram at 470 and 600°C [Figs 14(a) and 14(b)]. Fatemi and Weizer [115] have obtained ohmic contacts in the low 10m7 Rem’ range by annealing Ni (200 nm)/InP contacts (initial doping level of 1.7 x lo’* cmm3) at 400°C for 2-20 min.



Unlike Appelbaum et al. [116], they claim that N&P and not N&P, present at the semiconductor surface, is responsible for the low contact resistance. For longer annealing times, NiiP is replaced by N&P, which corresponds to a dramatic increase in contact resistance. Stremsdoerfer et al. [ill] have produced ohmic contacts through electroless deposition of Ni onto n-type InP (n = lOI cm-‘). Specific contact resistances on the order of lop6 I&m2 have been obtained over a wide temperature range (300-500°C). Unlike electroless deposited Pd/InP contacts, Ni reacts to form phosphides (Ni,P and N&P, which are metallic conductors with low bulk resistivities) instead of Ni-In compounds. Pt/InP metallizations. Reactions for the Pt/InP system are quite similar to the Ni/InP and Pd/InP systems. The main difference is that the reaction temperatures tend to be higher for Pt/InP [117], which is not surprising given the higher melting temperature of Pt relative to Pd and Ni. Ternary amorphous phase formation begins at about 325°C. This is followed by the formation of a crystalline ternary phase (Pt,InP) and a cubic phase, likely Pt,In with P in solution (35OC) [117]. Pt,InP does not form at the amorphous layer/semiconductor interface, but at the amorphous layer/metallization interface, and as such exhibits no preferred orientation with the InP substrate. This is quite different from both the Pd/InP and Ni,/InP systems where crystalline ternary phases nucleate and then grow with a preferred orientation at the semiconductor surface. A second ternary phase, which is denoted as Tl in Ref. [117], forms at 400°C followed by decomposition to binary phases at temperatures higher than 450°C. The final phases are PtIn, and PtP,, which agree with those expected from the bulk Pt-In-P phase diagram [118]. As with the Ni/InP system, the Pt/InP reactions are laterally non-uniform. Near-noble transition metal-based/InP contacts. Both Pd and Ni can be used alone in the formation of ohmic contacts to n-type InP, with Y, values as low as 10m6 Rcm2 being obtained for high initial doping levels (see the discussion in the previous section). Generally, however, Pd or Ni is used in combination with other elements. such as n- or p-type dopants and Au to form practical ohmic contacts on InP, often with lower initial doping levels (10’7-10’8 cm-‘). The main role of Pd and Ni in this case is to improve metallization adhesion to the semiconductor and to initiate decomposition of InP and subsequent outward diffusion of In. Platinum is most commonly used in combination with Ti, and in some cases Au, to form so-called “non-alloyed” contacts. Its role is not clearly defined, but appears to be as a barrier to both Au and In diffusion. Pd-based/ZnP contacts. Pd-based metallizations are utilized in ohmic contacts to both n- and p-type InP. One particular contact structure of interest to n-type InP is a Ge/Pd metallization. This structure is based on a successful, non-spiking contact to n-type GaAs [119-1241. Ohmic contact formation in the GeiPdiGaAs system is based on a solid phase regrowth mechanism [125]. According to this model, films of two elements, M and M’ (Pd and Ge for this specific case) are deposited sequentially on a compound semiconductor, AB. At low annealing temperatures, the M metal film adjacent to the semiconductor reacts with AB to form a ternary phase M,AB according to :






+ M,AB.

The second film (M’) is chosen such that it drives the decomposition of M,AB and forms a stable binary compound with M. In addition M’ should be a suitable dopant for AB : (xy)M’+

M,,AB -+ AB+xMMj.

AB also regrows (the regrown layer is often distinguishable from the substrate in TEM analysis because it is faulted), doped with M’, epitaxially on the AB substrate. Annealing temperatures are in the 300-500°C range. Specific contact resistances in the 10d6 Qcm’ range have been reported for GaAs doping levels of l-2 x 10” crne3 [120, 1221. The first published work on Ge/Pd contacts to n-type InP was reported in Ref. [ 1261. Only electrical results were reported, with contact resistances around 10m5 0cm2 for original InP doping levels of 5 x 10” cm-3. A more recent study was done on Ge/Pd contacts using backside SIMS to characterize the reaction products [127, 1281. In backside SIMS, depth profiling is done by sputtering from the substrate side instead of from the metallization side, which is more common. Layers of Pd (50 nm) and Ge (130 nm) were deposited (Pd/Ge atomic ratio < 1) onto InP, which was doped to an unspecified level. A Pd-InP ternary region was detected at the contact/semiconductor interface after 200°C annealing, which is consistent with previous findings for single-layer Pd films on InP [103, 1 lo]. In addition, Pd-Ge compound formation occurs (likely PdGe). For anneals at 325’C, Schwartz et al. [127, 1281 argue that some Pd-In-P decomposition has taken place, resulting in the formation of a lo-nm regrown layer of InP (Z 10 nm thick), doped with Ge. There is, however, no conclusive evidence for the formation of a regrown layer, due to the nature of SIMS. We recently examined Ge/Pd contacts to InP using TEM techniques [129, 1301. Slightly lower r, values were achieved (6 x lop6 Qcm’) with a lower original doping level (10” cme3). PdjGe ratios and sequences were varied in an effort to obtain the optimum ratio, from the standpoints of contact resistance, uniformity and stability. Pd and Ge were deposited as either bilayers (Pd adjacent to InP) or trilayers, i.e. Pd/Ge/Pd. The effect of layer sequencing was minimal. The Pd/Ge ratio does, however, affect contact resistance and microstructure. For Pd/Ge atomic ratios > 2. a minimum resistance of lo-’ Rem’ was obtained compared to 6 x 10m6 Rem* for Pd/Ge < 1. For Pd/Ge > 2, the reactions were essentially the same and are summarized in the equations below and the micrographs in Fig. 20. Ge+5Pd+InP

+ Pd,,InP+Pd,Ge

(as deposited)


-+ PdzInP+Pd,InP




4InP+ Pd,InP + 6Pd,InP(II)

2Pd,Ge + 31nP + Pd,InP(II)

+ 4PdIn + 2PdP, + 2[Ge]

3Pd,Ge + 21nP + 3PdGe + 2PdIn + PdP,


(425°C > 300 min).

Decomposition of Pd21nP(II) (400-425”C) corresponds to the onset of ohmic behaviour and attainment of minimal resistance values. Note that the reactions between Pd and InP are virtually the same as for single-layer Pd metallizations. In addition, Pd




Fig. 20. Bright field micrographs of Pd (25 nm)/Ge (5 nm)/Pd (10 nm) contacts to n-type InP: (a) as deposited ; (b) annealed at 250’ C for 10 s ; (c) annealed at 350°C for 20 s : (d) annealed at 425’C for 90 min ; (e) annealed at 425’C for 300 min [129]. forms binary compounds with Ge, i.e. Pd,Ge during deposition and PdGe at higher temperatures. Annealing at > 425°C results in further decomposition of InP and reversion to Schottky behaviour. Decreasing the PdjGe ratio (or increasing the relative amount of Ge) results in a proportionate decrease in the amount of Pd available for ternary phase formation, indicating that Pd-Ge phases are kinetically, if not thermodynamically, preferred over the ternary phases. This has the benefit of reducing the amount of InP consumed during ohmic contact formation. For a Pd/Ge atomic ratio of z 11, about 75 nm of InP is consumed, compared with 20 nm of InP for a Pd,/Ge ratio of 2.7. Decreasing the PdjGe ratio to below I .O significantly affects the reaction process. Pd,Ge formation is bypassed, so that PdGe is the first and only PddGe binary phase to form. The amount of InP consumed decreases even further (about 17 nm consumed) and contact morphology is more uniform (Fig. 21). In addition, no Pd-In or Pd-P binary phases form when Pd21nP(ll) decomposes; the final contact products are PdGe and a

thin layer of unreacted Ge on the surface. The amount of excess Ge depends on the Pd/Ge ratio. The whereabouts of the released In and P [from decomposition of Pd,lnP(ll)] has not been ascertained at this time. Some In and P were dissolved in the PdGe layer [Fig. 21(d)], but this does not account for the entire amount. If regrown InP were to form during Pd21nP(ll) decomposition, then a layer at least 10 nm thick would be expected. A regrown layer has not been detected at the metallization/semiconductor interface [Fig. 2 l(b)], which does not necessarily preclude its presence. It is possible that the layer is virtually defect-free. Long-term annealing (5 h at 425°C) does not result in any phase transformations ; however, the interface becomes rougher [Fig. 21(c)]. The contact resistance increases somewhat. barely maintaining ohmic behaviour after 5 h at 425°C. The increase in resistance is likely due to Ge diffusion away from the surface region into the semiconductor substrate. The addition of Au to Pd/Ge contacts has the effect of lowering contact resistance. In a study where 25-nm Pd, 50-nm Ge and 40-nm Au layers were deposited on n-type InP (n = lOI

0. G. IVEY

and PING


cm-‘) [129, 1311, minimum rc values of 2.5 x 10-O Qcm’ wcrc obtained. The optimum annealing temperature was in the 350375°C range. The reaction sequence can be summarized in the following equations : Pd + InP + PddIn-P Au + Ge + Pd + Pd-In-P + InP + PdGe+Au-GeePd-P+Au,,,In, Au-GeePd-P

+ InP + GeP + Au, Jn, + PdGe

(as deposited) (250-325‘C) (a325


The overall equations have not been balanced because of the uncertainty in the concentrations of the Pd-In-P and Au-Gee Pd-P phases. As with other other Pd containing metallizations. the role of Pd remains the same. Pd reacts with InP at low temperatures to form the amorphous PddInP phase. thereby improving contact adhesion and initiating InP decomposition. During annealing, however, the amorphous phase does not transform to crystalline ternary phases, but decomposes in the presence of InP. Au, Ge and Pd to form binary phases (PdGe and Au,&,) and a metastable quaternary phase (Au-Pd-GeP), which grows epitaxially on the InP. Ohmic behaviour occurs when the epitaxial phase decomposes-presumably some of the Ge released during this decomposition diffuses into the underlying InP. The overall balanced reaction for this system is :

and 3540 nm of 1nP is consumed in the reactions. Along with the lower contact resistance, relative to the Ge/Pd contacts, comes poorer contact stability and uniformity, in the form of spiking of a Au-In phase (Au,,,In,) into the InP (Fig. 22). Spiking becomes more pronounced at higher annealing temperatures (450°C) and GeP forms a continuous network. This corresponds to an increase in contact resistance and eventual reversion to Schottky behaviour. From the experimental observations of the two Pd-based contact systems described above (Ge/Pd and Au/Ge/Pd), some common characteristics are apparent : 1. Both the Au/Ge/Pd contacts and the GejPd contacts have strong metal,/InP adhesion due to the presence of Pd at the semiconductor interface, which reacts with InP to form an amorphous ternary phase. This indicates the importance of Pd in the metallization scheme in improving the wettability and adhesion at the InP surface. 2. The contact structures of both systems contain an epitaxial layer (quaternary phase in the Au/Ge/Pd/InP contacts, ternary phase in the Ge/Pd/InP contacts) in the early stages of annealing for each system. Neither phase is stable on InP, i.e. they are only transition phases. 3. In both systems, when the epitaxial phases decompose the contacts become ohmic, although the annealing temperatures required are different. The Au/Ge/Pd/InP contacts are ohmic in the 300-375°C range, while the Ge/Pd/InP contacts are ohmic in the 350-450°C range. (The temperature for the onset of ohmic behaviour decreases as the PdiGe ratio decreases.) There are also differences in terms of the electrical properties and contact morphology characteristics between the two





systems. The metallization-semiconductor interface in the Ge]Pd contacts is quite uniform, particularly for those contacts where the Pd/Ge ratio is slightly less than 1 (Fig. 21). Aucontaining contacts on the other hand are susceptible to spiking (Fig. 22). which is a negative consequence of gold’s affinity for In. The contact resistance for the Ge/Pd contacts is at least a factor of 2 higher than the AulGejPdiInP contacts. Thus, Au plays a role in lowering the contact resistance. Two possible mechanisms for ohmic contact formation have already been discussed, i.e. compounds with work functions less than the semiconductor may form at the metallizationj semiconductor interface resulting in a zero (or low) energy barrier, or dopant diffusion into the semiconductor could lead to the formation of a heavily doped InP surface layer. In the Au/Ge/Pd contacts. Au,&,, GeP and PdGe form at the metallization/InP interface. It can be inferred that Au,&, and GeP do not have lower work functions than InP, since both are present when the contacts are ohmic and in later stages of annealing when the contacts become Schottky barriers again. In the Ge/Pd contacts, PdIn, PdP,, PdzGe and PdGe can form during thermal reaction. PdIn and PdP, are not the compounds responsible for the onset of ohmic behavior in the Ge/Pd/InP contact system, for the same reasons that Au,&, and GeP are not in the AujGejPd system. In addition, PdIn and PdP, do not form in contacts with Pd:‘Ge ratios less than one. The Pd,Ge phase is in the outer layer of the contact when the contacts become ohmic, i.e. it is not in direct contact with InP, so that this phase is also ruled out. PdGe also cannot be responsible for ohmic behaviour, since in the Ge/Pd contacts with a Pd/Ge ratio of 11, it only forms at higher annealing temperatures when the contacts revert to Schottky behaviour. PdGe is therefore not likely to be responsible for ohmic behavior in Au/Ge/Pd/InP contacts either. From the above arguments, it can be concluded that ohmic contact formation in Au/Ge/Pd/InP contacts and Ge/Pd/InP contacts is not governed by the formation of zero energy (or low energy) barriers at the metallization/InP interface. It is our opinion that the onset of the ohmic behavior in the Au/Ge/Pd/InP contacts and the Ge/Pd/lnP is due to the formation of a shallow surface layer in the InP heavily doped by Ge diffusion. Gold enhances the doping process, thereby accounting for the reduction in contact resistance for Au/Ge/Pd contacts relative to Ge/Pd contacts. The following model is proposed to explain this effect. Au and Pd can react with In and P (as has already been discussed) forming various binary compounds in the AuiInP system and Pd/InP system, respectively. In the Au,/Ge/Pd/InP system, however, there is no Au-P compound formation because of competitive growth between the various compounds. Similarly, no Pd-P compounds are found for the AuiGejPdjInP system because Pd is restricted to PdGe phase formation due to the strongly preferred Pd-Ge reaction. The only phosphide formed during the entire annealing process is GeP. The net reactions between Au, Ge. In and P can be expressed by the following equation :

According to this equation, the volume ratio of GeP/Au,,In, would be 0.54 if the reaction process exactly followed the stoi-



and PING


chiometry in the equation. The reaction between Au and In starts at low temperatures (<25O’C), but the reaction between Ge and P begins at relatively high temperatures (> 3OO’C). The amount of GeP is much less than that of Au,&, in the contacts (Fig. 22), which indicates that the reaction rate for GeP formation is slower than that for Au,&, formation when the temperature is lower than 400°C. More In is consumed in the reactions than P, so that the P to In ratio in InP must be larger than 1 in the area near the metallization/InP interface. In other words, outward diffusion of In from InP leaves a large number of vacancies near the metallization/InP interface, providing sites for Ge doping. In the Ge/PdDnP contacts two possibilities exist. If there is excess Pd, i.e. the Pd/Ge ratio exceeds 1, then Pd reacts with both In and P [during Pd,InP(II) decomposition] to form PdIn and PdP, :

Fig. 21(a)-(c).




3Pd + 21nP + 2Pdln + PdP, The volume ratio of PdP,/PdIn is 0.36 if the amount of In and P. released during decomposition of the Pd-In-P epitaxial layer, is the same and all the released In and P react with Pd according to the above equation. In the Pd/Ge/Pd/InP contact system. the reaction between In and Pd and the reaction between P and Pd take place at the same temperature and the same time. Although the reaction rate is not known, it can be seen from the micrographs of specimens annealed at 425°C (Figs 20 and 22) that the difference between the actual volume ratio of PdP,/PdTn and the theoretical ratio of PdPJPdIn is less than that between the actual ratio of GePjAu,Jn, and the theoretical value of the GeP/Au,,,In, ratio. This implies that the number of In site vacancies in the Au/Ge/Pd/InP contacts is higher than in the Ge/Pd(InP contacts. Therefore. following the mechanism

Bright field micrographs of Ge (50 nm)/Pd (30 nm) contacts to n-type InP : (a) annealed 25O’C for 5 min; (b) annealed at 400°C for 5 s ; (c) annealed at 425°C for 5 h.






z-. *z


5 s 400

‘! In








Fig. 21(d). EDX spectrum from PdGe in (c), showing In and P in solution [130].

Fig. 22. TEM bright field micrographs from Au:Ge/Pd contacts annealed at (a) 400°C and (b) 450°C. Note, the “spiking” of Au,& into InP in (b) [129, 1311.

outlined above, the doping level of Ge in the InP substrates of the Au/Ge/Pd/InP contacts should be higher than that in the Ge/Pd/InP contacts. As a result, tunneling in the Au/Ge/Pd/InP contacts should be easier than in the Ge/Pd/InP contacts and, hence, the minimum contact resistance value for the Au/Gel PdiInP contacts should be lower than for the GejPd/lnP contacts. If there is excess Ge (Pd/Ge < l), the mechanism for ohmic contact formation is somewhat different. PdIn and PdP, do not form, since Pd is tied up in PdGe formation. This indicates that the driving force for PdGe formation exceeds that of PdIn and PdP>. Regrown InP may form (although conclusive evidence is lacking), which may be doped with some of the excess Ge, producing the low resistance tunneling contact. The doping process appears to be more efficient (lower r, values are achieved) compared to the Pd/Ge contacts where Pd/Ge > 1. The reversion to Schottky behaviour at higher annealing temperatures (or longer times) for the Au/Ge/Pd and Ge/Pd (Pd/Ge > 2) systems is likely due to further reaction with InP, which results in decomposition of the heavily doped n+ layer. For Ge/Pd contacts with Pd/Ge < I, ohmic behaviour is maintained for long annealing times and relatively high temperatures, because there are no further reactions after PdJnP(II) decomposes to form PdGe (Fig. 21). The contacts do eventually become Schottky barriers again, due to Ge diffusion away from the semiconductor surface. Other Ge/Pd type contacts include those with a Si capping layer instead of Au, e.g. Si/Pd/Ge (2 nm)/Pd/InP [128]. After annealing at 2OO’C, the contact consists of Pd$i, PdGe and Pd-In-P. SIMS depth profiles show that Ge is rejected by the Pd-In-P interfacial phase and resides in the Pd-Ge layer. Upon



and PING


further annealing (325°C for 30 min), the profile shows a welldefined Pd-SiLGe layer and diffusion of Si and Ge into the PdIn-P layer. The contact structure is Pd$i/Pd-Si&Ge/Pd-InP/InP. No electrical results are reported in the investigation. The details of the relationships between reactions and electrical properties in Pd-based,/InP contacts are not clear at this time. In addition, more work is required in order to understand the ohmic contact formation mechanism in Pd-based contacts. In addition to application on rz-type InP, Pd-based metallizations have been employed to p-type InP. For instance, Pd/Zn/PdiAu metallizations have been used to fabricate ohmic contacts on p-type InP [132]. Experimental results again show a strongcorrelation between the electrical results and the microstructural changes that occur during annealing. The role of Pd is the same as with the other Pd-containing contacts. The first reaction product is an amorphous layer, with a composition corresponding to PdJnP, at the Pd/InP interface. The amorphous layer crystallizes and grows epitaxially on InP by 35O‘C. Meanwhile, Au diffuses toward the PdzInP and InP, In diffuses out to the Au layer and Zn reacts with Pd forming a thin PdZn layer at the interface of Pd and PdJnP. The contact structure is Au(In.Pd,Zn)/Pd/PdZn/PdJnP(Au)/InP. and the contacts are still Schottky barriers at this stage. At 45O”C, Zn is liberated by PdZn decomposition and dissolves in the Au layer, in Pd,InP and presumably in InP, which leads to the onset of ohmic behavior with a minimum contact resistance value of 7 x 10e5 0cm2. Pd liberated during PdZn decomposition reacts with InP by diffusing through the Pd,InP layer, forming PdP, and PdJnP. The contact structure consists of Au(In,Pd,Zn)/PdPZ/Pd,InP(Au.Zn)/lnP and the morphology is laterally uniform. Further annealing results in formation of AuIn-Pd ternary layers at the contact surface and the PdPJInP interface, which results in a breakdown of the uniform contact structure and an increase in the contact resistance. To summarize this section, contact resistances for several Pd-

Table 6. Resistance lnitial Metallization 150 nm(Au-Ge)/30 nm Ni 125 nm(AwGe),I25 nm Ni 20 nm Ge/40 nm Au/l 40 nm Ni 40 nm Au:‘20 nm Ge/l40 nm Ni 50 nm Gel500 nm Au/5 nm Ni 60 nmjAu-<;e)/l5 nm Ni/40 nm Au 50 nm(Au+Ge)j75 nm NijlOO nm Au 25 nm Nil50 nm Gel40 nm Au 25 nm Nil50 nm Au/25 nm Ge:‘5 nm Ni 300 nm(Au-Ge)/70 nm Ni 40 nm Ni:300 nm(Au-Ge)/70 nm Ni 50 nm Pd/ll5 nm Ge 25 nm Pdi50 nm Ge/40 nm Au 10 nm Pdj5 nm Gei2.5 nm Pd 30 nm Pd/50 nm Ge 200 nm Ni 100 nm Ni,P Pd 10 nm Pd/5 nm Zn/20 tcontact

to p-type

nm Pd/60

nm Au+

InP substrate.

of near-noble doping (cm-‘) 8x10”

1.7 x 10IX 1.7x lOI lOI 2 x lo’x IO” j-9 x 10” 8 x 10’:

8x 10’:

5.3 x 10’” IO” 10” 1.7% 8 x lo’8 10’” 2 x 10IX


OHMIC CONTACTS based metallization> Table 6.

TO InP on w- and I>-InP substrates arc listed in

Gi-ba.sed/InP contacts. Au/Ge/Ni contacts are the most common contacts to n-type 1nP. based on similar successful ohmic contacts to n-type GaAs. The reactions between the metallization and InP substrate are complex, and no universally applicable current transport model for this contact system has emerged yet. Various arrangements of the multilayer contact have been studied (the metal listed last is the closest to the semiconductor), such as Ni/Au-Ge [133, 1341, NijAuiGe [77, 135, 136. 1411, Ni/Ge/Au [135], Au/Ni/AuGe [135, 137, 1381, Au/Ge/Ni [I 39, 1421 and AujNi/Ge/Ni [72, 1401. The result is a variation in the reaction products, morphology and electrical properties. From the results reported in the literature the characteristics of AuGeNiiInP contacts during processing can be summarized as follows: (1) Ge moves from the Au-Ge (or Ge) layer to the Ni during the initial reaction stage (< 250°C). (2) Ni, if separated from InP by other layer(s), diffuses toward the surface of the InP substrate. If Ni is in contact with InP. it reacts with InP at low temperatures (< 25O’C) to form a Nirich ternary amorphous phase. (3) In diffuses outward to the Au layer or the surface of the contacts. (4) Au migrates to the InP substrate. Several reaction products are reported in the literature : NiGe [77, 137, 139, 1421, Ni,P,8 (including NiP, N&P, N&P, N&P) [135, 139, 1401. Ni,lnP 1139, 1421. Au,In [139, 1421, AuZP, [136, 1411, AuGeP [1136], AuGeNiP [140] and Inz03 [140]. However, many of these can be viewed as being suspect due to the characterization techniques utilized. The reactions for one particular system [Au (40 nm)/Ge (50 nm)/Ni (25 nm)] are summarized in the scheme shown in Fig. 23 [142]. Ni, as in the single layer metallization, reacts with InP during deposition to form an amorphous ternary phase (10 nm thick) at the interface. In addition, some of the Ge reacts with Ni to form an amorphous Ge-Ni layer (lo-15 nm thick). The



Resistance @cm’) 0.09WI.16 Qmm 8 x 10-h IO-’ 10-S 1.5 x 10 6 2x10-’ 8x lo-: 1o-7 1o-4 1.2 x lom6 2.3 x 10m6 1.27 x lO-5 2.5x 10 ’ lo-5 6x 10m6 3 x 10-7 3 x lamb 5 x lo-’ 7 x loms

contacts Processing



360-630°C. 30 s 375’C. 3 min 4OO”C, < 20 min 4Oo”C, < 5 lnin 330-390°C. 15 min 430-45O’C, 2s 45O”C, 25 s 400°C. 10 s 45O”C, 1 min

133 134 135 135 136 137 138 139, 142 140 72 72 126 129. 131 129 130 115 116 111.112

35mOO’C. 2 min 400°C. 2 min 300°C 30%375°C 400425°C 400-45O’C 400°C Sputtered Electroless deposition 30~350°C 42&425”C, 30-70 s


D. G. IVEY and PlNG


as deposited





4oo” c

Fig. 23. Schematic representation of reactions that occur in the Au (40 nm)/Ge (50 nm):Ni (25 nm) system during annealing [ 1421.

remaining Ge is amorphous. Annealing at 250°C results in growth of the Ni-In-P ternary layer and crystallization of the Ni-Ge phase to NiGe. The remaining Ge is crystalline and intermixed with Au in the outer layer. In has begun to diffuse to the outer Au layer and Au has begun to diffuse towards InP, detected by EDX analysis. Ohmic behaviour is achieved at 325°C (rc = 8 x lo--‘ Qcm’). This corresponds to the decomposition of the Ni-In-P phase and the formation of N&P and Au,&, at the semiconductor surface. Annealing at 35&4OOC results in attainment of the minimum contact resistance (rc = 10M7 Rem’, for an initial doping level of 10” cm-“), which corresponds to further inward diffusion of Au and outward diffusion of In and formation of Au&r, (from Au&r,) and NiP (from Ni,P). The overall reaction for the contact is : 9Au + Ge + 5Ni + 41nP -+ Au,In,

+ 4NiP + NiGe.

The amount of InP consumed can be estimated from mass balance considerations and corresponds to around 60 nm. Only about 25 nm of the Ge (half the deposited thickness) is consumed in NiGe formation, leaving excess Ge in elemental form. As stated above, the change in electrical properties from Schottky to ohmic behaviour, appears to correlate with the decomposition of the ternary phase (NiJnP) and the formation of N&P and Au&,. Ge doping of InP may be a factor, as is often claimed, since Ge was detected in N&P, and N&P is in contact with InP in places, which would facilitate diffusion. Another possibility is the lowering of the barrier height at the interface, as proposed by several authors [116, 135-137, 140, 1411. Ohmic contacts have been achieved to n-type InP (doping level of 2 x 10’” cm-‘) by sputtering N&P (which is a metallic conductor with a bulk resistivity of 32 @cm) directly on the semiconductor surface [116] or through thermal reactions (at 400°C) between Ni films and InP (n = 1.7 x lOI cm-‘) [I 151. Further evidence pointing towards the lowering of the barrier height is provided by Fatemi and Weizer [ 1351, who compared Au/Ge/Ni contacts with Au/Ni contacts. Both systems exhibited ohmic behaviour after annealing to 400-C for a few minutes and had comparable resistances (Z IO-’ acm’). The drop in resistance appeared to correspond to the formation of nickel phosphides, and the resistance-temperature behaviour was virtually identical for both systems. It is likely that in the multi-layer structures (Au/Ge/Ni) both a barrier height lowering and a doping mechanism may be operative. In the system described in Ref. [142]. the contact

resistance decreases by almost three orders of magnitude (from 8 x low5 to lo-’ .Qcm2) when the annealing temperature is increased from 325 to 400°C. This phenomenon cannot be accounted by barrier height lowering alone, The onset of ohmic behaviour at 325°C may, in fact, be a result of a lower barrier height, due to interfacial contact between N&P and InP; however, at 350°C and above, Ni,P is no longer present, as it transforms to NIP. At these temperatures Ge diffusion is likely more important (NiGe is in contact with InP), accounting for the substantial decrease in contact resistance. In addition, for most of the systems reported in the literature, the initial doping level in InP was 10’” cmm3 or higher compared to the initial value of 10” crne3 in Ref. [142]. The InP in the systems reported in the literature has already been in effect inherently doped prior to annealing, providing some narrowing of the depletion zone or interface barrier. The role of Au in these contacts appears to be in helping to drive the reactions, in particular the decomposition of the NiIn-P ternary phase. Unfortunately, Au also contributes to interface roughness and contact instability (driving the outward diffusion of In). Contact resistances for various Au-Ni-Ge metallizations are given in Table 6, with values varying from lo-’ to 10m8 acm’. Pt-bused:‘InP contacts. The main application of Pt in InP-based semiconductors (including InGaAs and InGaAsP) is in socalled non-alloyed contacts. These usually consist of Ti and Pt layers. with or without a Au capping layer, deposited sequentially on the semiconductor [143-1581. Gold, which can be deposited either before or after contact annealing, is needed to permit wire bonding and solder bonding of the device to a suitable submount [144]. Deposition of Au prior to annealing permits deposition of all the metal layers in one sequence, without breaking the vacuum, although Au may reduce contact stability during subsequent annealing. The Pt and Ti layers serve to improve adhesion to the underlying semiconductor and act as a diffusion barrier for Au [ 1451. Good contacts have been obtained using this scheme to a number of semiconductors. with contact resistance values depending very much on the initial doping levels in the semiconductors. Several examples are shown in Table 7, including contact resistances, doping levels and optimum annealing conditions. The original doping levels in the semiconductors are generally quite high (Table 7) so that many of these contacts are ohmic without annealing or


D. G. IVEY Table

Doping level (cm-‘)

Semiconductor p-In,



n-In, ,Ga, ,As. Si-doped p-InGaAsP, Zn-doped



7. Pt!Ti-based

1.5 x 5x 5x 5x 2x lx 3x 2x 8x 5x 2-3 x 2-3 x 1x 5-8 x

10” 1o’x 10’” 10’s 10” 10’8 lOI 10” 10’s 10’s 10” 10” 10’s 1018

556 x lOI p-InAs,




1x10’s 5 x 1o1s 1 x 10’” 1.4 x lOI

and PING ohmic

JIAN contacts

: OHMIC to various





Metallizationt Pt (75 nm)/Ti (50 nm) Au (500 nm)/Pt (60 nm)/Ti Pt (60 nm)/Ti (50 nm) Pt (75 nm):‘Ti (75 nm) Pt (100 nm)/Pd (25 nm) Au (250 nm)/Pt (25 nm)/Ti Pt (75 nm)/Ti (50 nm) Pt (75 nm)/Ti (50 nm) Pt (75 nm)/Ti (50 nm) Au (150 nm),/Ti (30 nm) Au (150 nm)/Pt (50 nm)/Ti Au (500 nm)/Pt (60 nm)/Ti Pt (75 nm)iTi (50 nm) Pt (75 nm)/Ti

(50 nm)

Pt (60 nm)iTi Pt (60 nm)/Ti

(50 nm) (50 nm)

Pt (60 nm)/Ti (50 nm) Pt (100 nm)iTi (20 nm)

(50 nm)

(25 nm)

(30 nm) (50 nm)

Contact resistance (&m’)$


3.4 x IO-8 6 x IO-’ 4 x 10m6 9 x 10mb 1.2 x 10m6 (unannealed) 2 x 10d5 (4 min at 425’C) 5-6 x IO-’ (unannealed) 1.1 x 10mh 7.2 x 10m6 3.6 x 10mJ 1.6 x IO-’ (unannealed) 1.6x IO-” (unannealed) 8xlo-6 8 x 10-l 3 x lo-’ (unannealed) 1 x IOW 1.8 x 10m6 (unannealed) 14x 10.’ 1.1 x IO-” 3 x 10m9 999.9 x 10m8 1 x lo-6

143 144 151 145, 156 160 159 152 146 146 148 158 158 144 149 150 149 150 147, 153 151 153 147,153 154

j-First metal listed is on the outer surface of the contact. Metals were deposited by electron beam evaporation. $A11 contact resistance values listed are for rapid thermal processing anneals at 450°C for 30 s, unless otherwise

they require very little annealing to attain ohmic behaviourhence the term non-alloyed contacts. In reality, many of these contacts are annealed to optimize the electrical properties and therefore undergo fairly significant interfacial reactions (see discussion below), making the terminology “non-alloyed” something of a misnomer. Although many papers have been written describing the electrical and stress behaviour of Pt/Ti-based contacts, very little has been done until recently to characterize the microstructural changes that occur during annealing. The analysis that has been done is mainly by Katz et al. [143-151, 15331561. Most of their analysis has involved AES or Rutherford backscattering spectroscopy (RBS), with a limited amount of transmission electron microscopy (TEM). Their work is summarized in the following. In contacts to InGaAs, very little reaction occurs until 450°C (note, all of their contacts were annealed by rapid thermal processing or RTP), which also corresponds to a minimum in contact resistance [1433145, 151. 1561. At this temperature, there is significant interdiffusion between the Pt and Ti layers, resulting in a layer with a somewhat larger grain size relative to the as-deposited condition, i.e. 56 nm vs 11 nm [145, 1561. No further grain growth occurs at higher temperatures. There is also significant interaction between Ti and the underlying semiconductor, leading to Ti in-diffusion and In and As out-diffusion. The reacted zone at the interface consists of a “defective interfacial layer (or ‘deformed’ InGaAs) and solid state regrown second phase regions” [145, 1561. The “second phase region” has not been conclusively identified, although the authors speculate that it may be made up of any number of the following intermetallics : InAs, InTi,, Ga,Pt,, GaPt, Ti,Ga,,


In,Pt and GaPt, [143, 145, 1561. Similar behaviour is reported for Pt/Ti-based contacts to InGaAsP [ 1461, InP [ 1491 and InAs [153]. Again, only limited reaction occurs until about 45OC. whereupon Pt-Ti intermixing and reaction between Ti and the semiconductor take place. Most of the reaction products have not been conclusively identified. Detailed microstructural analysis was done by ourselves [ 1591 on contacts similar to p-type InGaAs (250 nm Au/25 nm Ptj 25 nm Ti). The microstructures were virtually the same as those reported by Katz et al. The onset of ohmic behaviour (3755 425°C) corresponds to Ti in-diffusion and In and As outdiffusion, resulting in significant microstructural changes at the Ti/InGaAs interface (Fig. 24). Some of the Ti reacts with As to form TiAs, while the In precipitates out as metallic In. A layer of “deformed” InGaAs, deficient in In and As, is left at the interface, producing a laterally nonuniform microstructure at the semiconductor surface. The Pt layer acts as a fairly good diffusion barrier (for anneals up to 45OC), preventing In outdiffusion into the Au layer and minimizing Au diffusion towards the semiconductor interface. The contact has very poor morphology, however, due in part to the low melting point In present at the metallization,/semiconductor interface. Katz et al. have proposed a phenomenological model to explain the contact resistance behaviour for Pt/Ti-based contacts [23, 145, 1561. For example, for PtiTi contacts to p-type InGaAs (p = 5 x 10” cm-‘) [23], the dominant carrier transport mode for as-deposited samples was found to be thermionic emission (calculated barrier height of 0.13 eV). For annealed samples, according to the model, carrier transport relies on both thermionic and field emission of carriers, with field emission


Fig. 24. TEM bright field micrographs taken from Au (250 nm)! Pt (25 nm)/Ti (25 nm) contact, annealed for 2 min at (a) 375°C and (b) 425°C [159].

dominating at the optimum annealing temperature. This corresponds to major changes in interfacial microstructure, as described above.

CONCLUSIONS The formation and characterization of ohmic contacts to InP, and to a lesser extent related ternary and quaternary semiconductors, have been discussed from a metallurgist’s point of view. Ohmic contacts are considered to be low resistance contacts (< 10m6 Rem’ for n-type InP and < lo-“ Qcm2 for p-type InP), formed through annealing of an appropriate metallization scheme. The mechanism of ohmic contact formation to InP is not well understood, although it is believed to be associated with doping of the underlying semiconductor, the formation of compounds which lower the barrier height at the metallizationisemiconductor interface or some combination of the two. Electrical requirements, i.e. primarily low contact resistances, have been considered in this paper; however, the discussion has focused primarily on important metallurgical aspects, such as contact stability, uniformity and semiconductor consumption. These issues become increasingly important as device sizes decrease. Microstructural characterization can be difficult, due to the small dimensions (total metal layer thickness is generally less than a few hundred nanometers) associated with these systems. Transmission electron microscopy (TEM) is shown to be a very powerful microstructural characterization



tool, providing high resolution imaging and high spatial resolution compositional and structural information in three dimensions. Most ohmic contact structures to InP are essentially empirical in nature, derived from similar, successful contacts to GaAs. The metallizations are, in general, multilayer systems, consisting of 3 or 4 elements in many cases. The number of metallic constitutents, coupled with 2 or more semiconductor components, results in an inherently unstable system. The function of each metallization component is not always clear, although intended roles include dopant species (e.g. Ge for n-type doping and Zn for p-type doping), adhesion promoters (e.g. Pd and Ni), protective capping layers (e.g. Au) and diffusion barriers (e.g. Ti and Pt). Gold-based contacts generally produce the lowest resistance contacts, but suffer from instability and lateral non-uniformity (or “spiking”), thereby compromising contact reliability. Gold has a high affinity for In. which drives the decomposition of InP and subsequent outward diffusion of In. This effect can be minimized to some degree by reducing the Au layer thickness or through the introduction of a diffusion barrier such as Pt or Ti. A more effective way of improving contact stability is to eliminate Au altogether and to reduce the number of metallization components. In addition, semiconductor consumption can be minimized by optimizing composition ratios. This was demonstrated for a Pd/Ge contact to n-type InP. Unfortunately, contact stability is often achieved at the expense of contact resistance, although this effect can be offset somewhat by increasing the initial doping level in the semiconductor. Acknowledgements-The authors would like to thank the Natural Science and Engineering Research Council (NSERC) of Canada and Bell Northern Research (BNR) Ltd for providing financial support for much of the experimental work presented in this paper.

REFERENCES 1. T. C. Eschbich, R. D. Carroll, R. N. Sacks and W. J. Tanski, IEEE. Trans. Elecrron. Devices ED-36. 1213 (1989). 2. J. A. Del Alamo and J. Mizutani, Solid-& Electr. 31, 1635 (1988). 3. P. E. Hallali, P. Blanconnier, L. Bricard and J. C. Renaud, J. Phys. CoNoq. 49, 453 (1988). 4. J. B. Boos and W. Kruppa, J. Vat. Sci. Technol. B 7, 502 (1989). 5. H. Schumacher, J. R. Hayes, R. Bhat and M. Koza. Int. Electron. Druices Meet, Technical Diqesi. IEEE, Washington, D.C. (1987). 6. A. Yoshida. H. Tamural T. Fujii and S. ‘Hasuo, Extended Abstracts of 1987 Int. Suoerconductivitv Electr. Conf. [SEC, 368 (1987). _I

7. H. Tenkin, R. A. Logan, R. K. Karlicek Jr, K. E. Strege, J. P. Blaha and P. M. Gabla.

Appl. Phys. Lett. 53, 1156 (1988).

8. R. Kaumans. N. Grote, H. G. Bach and F. Fidorra, Ins?. Phys. Conf. Ser. 91, 501 (1987). 9. M. Fukada. 0. Fuiita and S. Uetara. J. Liahtwaue Tech&. 6, 1808 (1988): ’ 10. D. L. Meier and D. R. Schroeder, IEEE Trans. Elrctr. Devices ED-31, 647 (1984). 11. T. S. Kuan, P. E. Batson, T. N. Jackson, H. Rupprecht and E. L. Wilkie, J. Appl. Phys. 54. 6952 (1983). 12. R. A. Bruce and G. R. Piercv. Solid-State Electr. 30. 729 (1987). 13. T. Sands, V. G. Keramidas, A. J. Yu, K. M. Yu, R. Gronsky and J. Washburn, J. Mater. Res. 2, 262 (1987). 14. T. Sands. Mater. Sci. Engng Bl, 289 (1989). I




and PING


15. C. Kittel, Introduction to Solid State Physics, p. 183. John Wiley and Sons. Inc.. New York (19X6). 16. J. W. Mayer and S. S. Lau, Electronic Materials Science: For Integrated Circuits in Si and GaAs. Macmillan. New York (1990). 17. M. L. Cohen and J. R. Chelikowsky, Electronic Structure and Optical Properties qj’ Semiconductors. Springer-Verlag, Berlin (1988). of Semiconductor Deaices. McGrawIX. D. S. Yang, Fundamentals Hill, New York (197X). 19. B. G. Streetman, So/id State Electronic Devices. Prentice-Hall. Inc., Englewood Cliffs, NJ (1972). 20. S. Adachi and J. Brice (eds). Proaerties of‘lndium Phosohide. The Institute of Electrical Engineers’London(l991). 21. J. M. Golio and R. J. Trew, IEEE Trans. Electron. Devices ED-30, 1411 (1983). 22. H. Morkoc, J. T. Andrewsand S. B. Hyder, IEEE. Trans. Electron. Deoices ED-26, 238 (1979). 23. A. Katz. In Iridium Phosphide and Related Materials : Processing, Technology, and Devices, p. 307 (edited by A. Katz). Artech House, Boston (1992). 24. G. Y. Robinson. In Physics and Chemistry qfIII&V Compounds: Semiconductor Interfaces (edited by Carl W. Wilmsen). Plenum Press, London (19X5). 25. E. H. Rhoderich and R. H. Williams, Metal-Semiconductor Contacts. 2nd edn. Clarendon Press, Oxford (1988). 26. S. M. Sze, Physics ofSemiconductor Droices. John Wiley and Sons Inc., New York (1969). 21. T. C. Shen. G. B. Gao and H. Morkoc, J. Vat. Sci. Technol. B. 10, 2113 (1992). 28. D. G. Ivey and G. R. Piercy, J. Elect. Microsc. Tech. 8,233 (1988). 29. J. C. Bravman and R. Sinclair. J. Elect. Microsc. Tech. 1, 53 (1984). 30. N. G. Chew and A. G. Cullis, Appl. PhFs. Lett. 44, 142 (1984). 31. M. H. Loretto, Electron Beam Analysu of Materials. Chapman and Hall, Ltd ., London (1984). 32. J. M. Vandenberg and E. Kinsbron. Thin Solid Films 65, 259 (1980). 33. T. Yoshiie. C. L. Batter and A. G. Milnes, Thin Solid Films 111, 149 (1984). 34. A. J. Barcz, Mat. Res. Sot. Symp. Proc. 260, 3 17 (1992). 35. C. H. Mueller, P. H. Holloway and R. G. Connell, Mat. Res. Sot. Symp. Proc. 260.481 (1992). 36. Q. H. Hu. A. Kvist and H. 0. And&, Insr. Phys. Con/. Ser. 117, 91 (1991). 31. V. G. Weizer and N. S. Fatemi, J. Appl. Phys. 69. X253 (1991). 38. B. P&z, R. Veresegyhazy. G. Radnoczi, A. Barna, 1. Mojzes, 0. Geszti, and Gy. Vincze. J. Appl. PhJss. 70, 332 (1991). 39. B. P&z, G. Radnoczi, A. Barna, R. Veresegyhizy and I. Mojzes, Inst. Phys. ConJ Ser. 117, 275 (1991). 40. A. Hiraki, K. Shuto, S. Kim, W. Kammura and M. Iwami, Appl. Phys. Letr. 31, 611 (1977). 41. C. T. Tsai and R. S. Williams, J. Muter. Res. 1, 352 (19X6). 42. P. W. Chye, I. Lindau, P. Pianetta, C. M. Garner, C. Y. Su, and W. E. Spicer, Phys. R~L$. B 18, 5545 (1978). 43. J. M. Vandenberg, H. Temkin. R. A. Hamm and M. A. DiGiuseppe, J. Appl. Phys. 53, 7385 (1982). 44. J. Vandenberg, H. Temkin. R. A. Hamm and M. A. DiGuiseppe, Thin Solid Films 104, 419 (19X3). 45. H. Okamoto and T. B. Massalski, Binary* Alloy Phase Diagrams, D. 268 (edited bv T. B. Massalski). ASM. Metals Park. OH i1986). . 46. S. E. R. Hiscocks and W. H. Rothery. Proc. Res. Sot. London Ser. A 282, 318 (1964). 41. P. Villars and L. D. Calvert, Pearson’s Handbook of Crystullographic Data for Intermetallic Phases. American Society for Metals, Metals Park, OH (1986). 4X. M. Z. Jandali, T. Rajasekharan and K. Schutbert. Z. Metallkd. 73,463 (1982). 49. A. Hiraki. S. Kim, W. Kammura and M. Iwami. Surfuce Sci. 86, 706 (1979). 50. L. J. Brillson, C. F. Brucker, A. D. Katnani, N. G. Stoffel and G. Margaritondo. J. Vat. Sci. Technol. 19, 661 (1981). 51. C. T. Tsai and R. S. Williams. J. Mater. Res. 1, 820 (1986).




52. A. Piotrowska, P. Auvray, A. Guivarch, G. Pelous and P. Henoc, J. Appl. Phys. 52, 5112 (19X1). 53. 0. Wada, J. Appl. Phys.57, 1901 (1985). 54. A. Barna and B. P&z, J. Electron Microsc. Tech. 18, 325 (1991). 55. N. S. Fatemi and V. G. Weizer, J. Appl. Phys. 67, 1934 (1990). 56. V. G. Weizer and N. S. Fatemi, J. Appl. Phys. 68.2275 (1990). 57. N. S. Fatemi and V. G. Weizer, J. Appl. Phys. 65, 2111 (1989). 58. N. S. Fatemi and V. G. Weizer, Mat. Res. Sot. Symp. Proc. 181, 417 (1990). 59. N. S. Fatemi and V. G. Weizer, Appl. Phys. Left. 57, 500 (1990). 60. N. Szvdlo and J. Olivier, J. Ap~l. Phys. 50, 1445 (1979). 61. H. B.-Kim, A. F. Lovas, G.-G. Sweeney, T. M. S. Heng, Inst. Phys. Conf. Ser. B33, 145 (1977). 62. J. H. Pugh and R. S. Williams, J. Mater. Res. 1, 343 (1986). 63. I. Mojzes, D. Szigethy, and R. Veresegyhazy, Electron. Lett. 19. 117 (1983). 64. E. H. Rhoderick and R. H. Williams, Metal-Semiconductor Contacts, 2nd edn, pp. 745. Clarendon Press, Oxford (1988). 65. B. Tuck, K. T. Ip and L. F. Eastman, Thin Solid Films 55, 41 (1978). 66. A. Katz and S. J. Pearton, J. Vuc. Sci. Technol. B 9, 178 (1991). 67. S. C. Binari and J. B. Boos, Electron. Lett. 25, 1207 (19X9). Electr. 24,907 (1981). 68. P. A. Barnes and R. S. Williams. Solid-State 69. J. Dunn and G. B. Stringfellow; J. Electr. Mat. 19, LI (1990). A. Piotrowska, A. Barcz, J. Adamczewska and A. 70. E. Kaminska, Turos, Solid-State Electr. 29, 279 (1986). 71. 0. Oparaku, C. L. Dargan, N. M. Pearsall and R. Hill, Semicond. Sci. Technol. 5, 65 (1990). 72. A. Piotrowska, E. Kaminska, A. Barcz, J. Adamczewska and A. Turos. Thin Solid Films 130.231 (1985). R. J. McCoy, V. G.‘Keramidas and W. A. Bonner, 73. H. Temkin, Appl. Phys. Lett. 36, 444 (1980). J. Vat. Sci. Technol. A 1, 466 74. A. J. Valois and G. Y. Robinson, (19X3). K. Vogel and J. Zelinka, Semicond. Sci. Technol. 3, 75. V. Malina, 1015 (1988). 76. D. G. Ivey, P. Jian and R. Bruce, Thin Solid Films 190,217 (1990). A. Waseem and G. Y. Robinson, Thin Solid Films 77. L. P. Erickson, 64.421 (1979). 78. D. G. Ivey. R. Bruce and G. R. Piercy, J. Electr. Mat. 17, 373 (1988). 79. V. G. Keramidas, Thin Solid Films 96, 347 (1982). and E. Garmire, J. Appl. Phys. 61, 808 (1987) 80. T. C. Hasenberg Xl. J. B. Boos and W. Kruppa, Solid-State Electr. 31. 127 (198X). J. Appl. Phys. 73, 2353 (1993) 82. V. G. Weizer and N. S..Fatemi, X3. J. S. K. Mills and D. L. Kirk. Thin Solid Films 55, 149 (197X). 84. D. L. Kirk, J. S. K. Mills and J. E. Pattison, J. Phys. D 12, 1995 (1979). X.5. C. J. Jones and D. L. Kirk, J. PhJ*s. D 12, 941 (1979). 86. D. L. Kirk and J. S. K. Mills, Thin Solid Films 67, L29 (1980). 87. A. Saidane and D. L. Kirk, J. Phys. D 18, 1609 (1985). xx. J. J. Kelly, J. M. G. Rikken, J. W. M. Jacobs and A. Valster, J. Var. Sri. Technol. B. 6,4X (1988). 89. E. Kuphal, Solid-State Electr. 24, 69 (1981). 90. P. Auvray, A. Guivarc’h, H. L’haridon, J. P. Mercier and P. Henoc. Thin Solid Films 127, 39 (1985). 91. A. J. Valois and G. Y. Robinson, Solid-State Electr. 25, 973 (1982). 92. V. G. Weizer and N. S. Fatemi, Appl. Phys. Lett. 62. 2731 (1993). A. K. Chin. F. Ermanis, M. A. DiGiuseppe. J. A. 93. I. Camlibel, Lourenco and W. A. Bonner, J. Electrochem. Sot. 129, 2585 (1982). 94. S. E. Mohney and Y. A. Chang, Mat. Res. Sot. S.vmp. Proc. 260, 519 (1992). 95. S. E. Mohney and Y. A. Chang. J. Mater. Res. 7,955 (1992). 96. M. Andersson-Siiderberg, Mat. Res. Sot. Symp. Proc. 204, 231 (1991). 97. W. B. Pearson, A Handbook of Lattice Spacing and Structures of Metals. Pergamon Press, London (1967). 98. H. Okamoto, ASM Handbook : Volume 3, Alloy Phase Diagrams, p. 256. ASM International, Metals Park, OH (1992). 99. H. Okamoto, ASM Handbook : Volume 3, Alloy Phase Diagrams. p. 330. ASM International, Metals Park, OH (1992).

D. G. IVEY 100. 101. 102. 103. 104. 105. 106. 107. 108.

109. 110. 111. 112. 113.

114. 115. 116. 117. 118. 119. 120.

121. 122.



125. 126. 127.


129. 130. 131.

and PING


S. Rundqvist, ,I’,ture 185, 3 1 (1960). L. 0. Gullman. J. Less-Common Met& 11, 157 (1966). Y. Andersson, Actu Gem. Stand. A 31, 354 (1977). D. G. Ivey, P. Jian and R. Bruce. J. Electron. Mater. 21. 831 (1992). M. F. Singleton and P. Nash, Binary A//oJ, Phase Diagrams, p. 1368 (edited by T. B. Massalski). ASM, Metals Park, OH (1986). P. Nash, Bimwy Alloy Phase Diagrams. p. 1738 (edited by 7. B. Massalski). ASM, Metals Park. OH (1986). B. Bhattacharya and D. B. Masson, Muter. Sci. Engnq 22, 133 (1976). M. Andersson-Sliderberg and Y. Andersson. J. Solid-State Chem. 85,315 (1990). D. G. Ivey, D. Wang and P. Jian, Experimmtul Methods of Phase Diagram Determinution. p. 151 (edited by J. E. Mortal, R. S. Schiffman and S. M. Merchant). TMS Publications. Warrendale PA (1994). D. G. Ivey. L. Zhang and P. Jian. J. Mater. Sci.: Mater. in Electron. 2, 21 (1991). R. Caron-Popowich, J. Washburn, T. Sands and A. S. Kaplan, J. Appl. Phys. 64, 4909 (1988). G. Stremsdoerfer, Y. Wang. J. R. Martin and E. Souteyrand, Mat. Res. Sot. Symp. Proc. 260. 543 (1992). G. Stremsdoerfer, C. Calais, J. R. Martin, P. Clechet and D. Nguyen, J. Electrochem. Sot. 137, 835 (1990). A. Appelbaum, P. M. Thomas and P. A. Barnes. In Semiconductor-bused Heterostructures: inte+cial Structure and StubUity, p. 409 (edited by M. L. Green. E. E. Baglin, G. Y. Chin. H. W. Deckman, W. Mayo and D. Narashinham). The Metallurgical Society, Inc., Warrendale. PA (1986). T. Sands, C. C. Chang, A. S. Kaplan, V. G. Keramidas, K. M. Krishnan and .I. Washburn, Appl. Phys. Lett. 50, 1346 (1987). N. S. Fatemi and V. G. Weizer, J. Appl. Phys. 73, 289 (1993). A. Appelbaum. M. Robbins and F. Schrey, IEEE Trans. Electr. Devices ED-34. 1026 (1987). S. E. Mohney and Y. A. Chang, J. Appl. Phys. 74.4403 (1993). C. F. Lin, S. E. Mohney and Y. A. Chang, J. Appl. Ph~~s. 74.4398 (1993). L. C. Wang, S. S. Lau, E. K. Hsieh and J. R. Velebir, Appl. Phys. Lett. 54, 2677 (1989). C. J. Palmstrsm, S. A. Schwara. E. Yablonovitch, J. P. Harbison, C. L. Schwartz, L. T. Florez, T. J. Gmitter, E. D. Marshall and S. S. Lau, J. Appl. PhJs. 67, 334 (1990). E. D. Marshall, W. X. Chen. C. S. Wu. S. S. Lau and T. F. Kuech. Appl. Phys. Lett. 47. 298 (1985). E. D. Marshall, B. Zhane. L. C. Wane. P. F. Jiao. W. X. Chen. T. Sawada. S. S. Lau, K.-L. Kavanagh-and T. F. Kuech, J. Appl. Phys. 62,942 (1987). A. Paccagnella, L. C. Wang, C. Canali, G. Castellaneta. M. Dapor. G. Donzelli. E. Zanoni and S. S. Lau. Thin Solid Films 187,9 (1990). S. A. Schwarz, C. J. Palmstrom. C. L. Schwartz. T. Sands, L. G. Shanthdrama, J. P. Harbison. L. T. Florez, E. D. Marshall, C. C. Han, S. S. Lau. L. H. Allen and J. W. Mayer. J. VW. Sci. Technol. A 8.2079 (1990). T. Sands, E. D. Marshall and L. C. Wang. J. Muter. Res 3. 914 (1988). W. X. Chen, S. C. Hsueh, P. K. L. Yu and S. S. Lau. IEEE Electron. Device Lett. EDL-7. 471 (1986). S. A. Schwarz, M. A. A. Pudensi, T. Sands, T. J. Gmitter, R. Bhat, M. Koza. L. C. Wang and S. S. Lau. Appl. Phys. Lett. 60, 1123 (1992). S. A. Schwarz. T. Sands. R. Bhat, M. Koza, M. A. A. Pudensi, L. C. Wang and S. S. Lau, Mat. Res. Sot. Qmp. Proc. 260, 525 (1992). P. Jian. D. G. Ivey, R. Bruce and G. Knight. J. Ekctr. Mater. 23, Y53 (lYY4). P. Jian, D. Ivey and R. Bruce, unpublished results (1994). P. Jian. D. G. lvey. R. Bruce and G. Knight, Mar. Res. Sot. Symp. Proc. 260, 531 (1992).





132. D. G. Ivey, P. Jian, L. Wan, R. Bruce. S. Either and C. Blaauw. J. Electr. Mater, 20, 237 (1991). Solid-State Efectr. 31, 1635 133. J. A. Del Alamo and T. Mizutani, (1988). D. Gutierrez and 0. Aina, Solid-Stute 134. K. P. Pande, E. Martin. Eke.. 30. 253 (1987). 135. N. S. Fatemi and V. C. Weizer, Mat. Res. Sot. Symp. Proc. 260, 537 (1992). 136. D. A. Anderson, R. J. Graham and J. W. Steeds, Semicond. Sci. Technol. 3, 63 (1988). 137. G. Bahir, J. L. Merz, J. R. Abelson and T. W. Sigmon, 1. Electr. Mater. 16, 257 (1987). 138. H. Morkoc, T. J. Drummond and C. M. Stanchak, IEEE Truns. Elecfr. Derices ED-28, 1 (198 1). 139. D. G. Ivey, D. Wang and D. Yang, Microscopical Sot. Can. 20th Ann. Meet. 3-5 June 1993. Toronto, Ont. 140. M. F. J. O’Keefe, R. E. Miles and M. J. Howes, Iridium Phosphide and Related Mater. ,for Adv. Electr. Optical Devices 1144, 361 (1989). 141. R. J. Gr ,..,m, S. Myhajlenko and J. W. Steeds, J. Appl. Phys. 57. 1311 (1985). 142. D. G. Ivey, D. Wang. D. Yang, R. Bruce and G. Knight, J. Electr. Mater. 23.441 (1994). 143. A. Katz, W. C. Dautremont-Smith, S. N. G. Chu. P. M. Thomas. L. A. Koszi, J. W. Lee, V. G. Riggs, R. L. Brown, S. G. Napholtz and J. L. Zilko, Appl. Phys. Lett. 54, 2306 (1989). 144. A. Katz, B. E. Weir and W. C. Dautremont-Smith, J. Appl. Phys. 68. 1123 (1990). 145. S. N. G. Chu, A. Katz, T. Boone, P. M. Thomas, V. G. Riggs, W. C. Dautremont-Smith and W. D. Johnson Jr. J. Appl. Phys. 67, 3754 (1990). 146. A. Katz, P. M. Thomas, S. N. G. Chu, W. C. DautremontSmith. R. G. Sobers and S. G. Napholtz. J. Appl. Phys. 67. 884 (1990). 147. A. Katz, S. N. G. Chu, B. E. Weir, W. Savin. D. W. Harris, W. C. Dautremont-Smith. T. Tanbun-Ek and R. A. Logan, J. Vat. Sci. Technol. B 8, 1125 (1990). 148. A. Katz, W. C. Dautremont-Smith. P. M. Thomas, L. A. Koszi, J. W. Lee. V. G. Riggs, R. L. Brown. J. L. Zilko and A. Lahav, J. Appl. Phys. 65, 43 I9 (1990). 149. A. Katz, B. E. Weir, S. N. G. Chu, P. M. Thomas, M. Soler, T. Boone and W. C. Dautremont-Smith. J. Appl. PhyS. 67, 3872 (1990). 150. W. C. Dautremont-Smith, P. A. Barnes and J. W. Stayt Jr, J. Vat. Sci. Technol. B 2, 620 (1984). 151. A. Katz. S. N. G. Chu. B. E. Weir. C. R. Abernathv, W. S. Hobson, S. J. Pearton and W. Savin, IEEE Truns. Electr. Devices 39, 184 (1992). 152. S. Kuroda, N. Harad, T. Katakami and T. Mimura, IEEE Electr. Devices Lett. EDL-8, 389 (1987). 153. A. Katz, S. N. G. Chu, B. E. Weir, W. C. Dautremont-Smith, R. A. Logan. T. Tabun-Ek. W. Savin and D. W. Harris. J. At&. ._ Phys. &,4141 (1990). 154. A. Katz and W. C. Dautremont-Smith. J. Appl. Phys. 67, 6237 (1990). 155. W. Savin. B. E. Weir. A. Katz, S. N. G. Chu, S. Nakahara and D. W. Harris, iwut. Res. Sot. Symp. Proc. 181. 227 (1990). 156. S. N. G. Chu. A. Katz. T. Boone. P. M. Thomas. V. G. Riggs, W. C. Dautremont-Smith and W. D. Johnston Jr, Mat. R&y Sot. Qmp. Proc. 181, 389 (1990). 157. T. S. Kalkur. P. D. Wright. S. K. Ko, Y. C. Lu, L. Casas and K. A. Jones, Mat. Res. Sot. Symp. Proc. 260, 549 (1992). 158. V. Malina, E. Hajkova, Z. Zelinka, M. Dapor and V. Micheli, Thin Solid Films 223. 146 (1993). 159. D. G. Ivey, P. Jian, R. Bruce and G. Knight J. Mater. Sci. :Mafer. in Electronics (in press). 160. P. Ressel, K. Vogel, D. Fritssche and K. Mause. Electron. Let!. 28. 2237 (1992).