A review of the properties of aluminium alloy films used during silicon device fabrication

Vacuum/volume 42/number 12jpages 719 to 72911991 Printed in Great Britain

0042-207x/91 $3.00+.00 @ 1991 Pergamon Press plc

Invited Review A review of the properties of aluminium alloy films used during silicon device fabrication R J Wilson and B L Weiss, Department of Electronic and Electrical Engineering, University of Surrey, Guildford, Surrey GU2 5XH, UK

received for publication 17 December 1990

Aluminium alloy films have found widespread application during the metallisation stage of silicon device fabrication. The properties of these films are reviewed and links with the film microstructure resulting from a given set of deposition conditions are identified. Possible applications arising from the data presented to deposition control and film monitoring are discussed.

Introduction

Aluminium alloy films arc frequently used during the mctallisation stage ofdiscrctc or integrated circuit fabrication. They have evolved from the initial use of pure aluminium fihns which satistied many of the requircmcnts of the metallisation layer needed during silicon device fabrication ‘, The scaling of devices to improve performance led to the USC of AILSi alloys to suppress ;I failure mode associated with dissolution of silicon from contacts to silicon and subsequent penetration of Al into the contact,

Icading to a short circuit across the junction’.‘. The need to deposit an alloy of reproducible composition was one of the major factors leading to the preferred method of deposition changing from evaporation to dc magnetron sputtering’-“. Other additions to the aluminium alloy film. such as Cu. Ti and Pd, have been used to suppress other dcvicc failure modes and are discussed where relevant in the following sections.

A diagram illustrating the factors leading to a dcpositcd film with a given set of properties is shown in Figure I, This rcvicw

Vacuum Environment

Machine

Design

~ + Substrate

attempts to illustrate links between the separate factors leading to ;I final set of propcrtics. Much work neglects one or more 01 the links in the process, for cxamplc reporting the effect of process paramctcrs on fihn propcrtics without ;m understanding of the microstructure responsible for the film properties. Only with a knowledge of the complctc picture can a deposition process bc successfully controlled to give a film of the required quality. Currently the area of measurement of the fundamental paramctcrs has received little attention. Only when this is understood will in .situ monitoring. as opposed to the current (1.~ situ monitoring of film properties. become a viable technique for process control.

Finally the review discusses areas where the data from the literature suggest possibilities for improved deposition monitoring and control. which may lead to an improvement in equip- ment up-time and device reliability.

The following sections will discuss the electromigration resistance, specular rcflcctivity. rcsistivity, stress and step covcragc of deposited aluminium alloy films. Both clcctromigration and stress induced migration are becoming increasingly important as device dimensions arc reduced. Resistivity and specular rctlec- tivity arc often used as parameters to quantify the deposition process. whilst step covcragc is important since it can provide sites of prcfcrcntial failure by electromigration or stress induced migration.

Electromigration resistance

Elcctromigration is the current induced mass transport of a material. When a metal has a direct current flowing through it two forces act on a thermally activated metal ion. There is the force due to the applied electric held tending to move the ion to the negative terminal and the force due to momentum exchange bctwccn conduction electrons and the ion which tends to move the ion to the positive terminal. The force due to the cxchangc of momentum between conduction clcctrons (the clcctron wind) and the ion dominates since the conduction clcctrons are effective in screening the clcctric ficld from the ion’. Elcctromigration failure of conductor lines occurs at sites of metal flux divergence. which may occur because ofstructural inhomogeneitics. changes in the diffusion cocfhcient or temperature gradients (Figure 2).

719

RJ Wilson a&B L Weiss Alumlnltim alloy film propertles

J, J, Jo = J > No Damage

-- J. > J 2 Hillock

J.<J EVoid

I, AI, Other Cond,,,ons Reman Constant Then A Flux Discontinuity Occurs Where The Grain S’ze Al!ers

The Small Grain Section Can Transport More Materlai Than The Large GraIned Section And A Hillock %rms At AA

R./ Wilson andB L Weiss: Aluminium alloy film propertles

Table I. Studies showing the dependence of electromigration performance on the deposited film’s structure

Reference

I9 (196X)

5 (1969)

20 ( 1970)

21 (1981)

Film struclure:test result

Si substrate covcred over half its surface with 600 8, of thermal

oxide Orientation (I I I) on Si mixed on oxide for a deposited Al film Grain size equal on Si and oxide Si-oxide step was a site of electromigration failure

Cold Al evaporation (< 100 C), grain size < I .2 jtrn

Ilot Al evaporation ( - 400 C). grain size > 8 ltrn Activation energy : hot evaporation 0.84 cV : cold evaporation 0.48 eV /,,, (hot) > I,,, (cold)

Al cwporation 200 C grain size 1.2 Am mlxcd orientation

Al evaporatwn 500 C grain size 7.8 /cm (I I I) orientation

Activation energy : small grain size 0.51 cV ; large grain si7e 0.73 cV 80% of Ihilure at abrupt grain size changes

Al 0.5% Cu. Al 2”& Si and Al I”/o~ SiLO.25% Cu deposited using evaporation and sputterins No activation energy measurement

Grain size and film orientation measured

Films annealed prior to measurements

Conclusion

The change in film orientation from random lo (I I I) provided a barrier to clectrotransport

The increase in Ti,, for large grain films was because of II reduction in grain boundaries

The activation energy for clcctrotransport is 1~ function of the grain boundary structure Large grain equiaxed. ordered films are required for large I,,,

/,,, % microstructural parameter (A)

/r = (s’a’) log,,, (I, , ,:/w,J ’ s = median grain size (r = grain size standard deviation for a log-normal distribution I,,; = X-ray diffraction intensity from (_Y~Y) measured using an X-ray diffractometer

The link between electromigration performance and film structure highlights the riced to investigate systematically the effect of deposition parameters on film structure and therefore on clec- tromigration rcsistancc. Many factors can affect the final film structure and there has been much effort regarding the properties of evaporated Al and Al alloy films with less attention being paid to sputtered films. For evaporated films, the film orientation has been shown to be a function of deposition temperature’“.“, substrate type”.“. film thickness’“, residual gas pressure” I’ and evaporation rate”. Table 2. summarises some data regarding the orientation of sputtcrcd Al and Al alloy films, which dcm- onstrates the effect of deposition temperature”,‘“,“, residual gas pressure “, applied dc bias” 15, target design “.j4 and annealing conditions” on the fihn orientation. A comprehensive survey of the effects of deposition temperature, substrate type, number of substrate rotations past the target and substrate position in a batch mode system on film structure has been given by Wilson and Weissi’. There is a tendency for aluminium to form films with a (I I I) orientation. with the change in the strength of this orientation and the appearance of other orientations being linked to factors preventing grain growth, such as residual gas incorporation’~.“‘“‘“. sputtering gas incorporation’” and nucleation of grains away from the film~substrdte intcrfacc during bias sputtering”. Queirolo c’t (11’” have shown that the distribution of planes with rcspcct to the substrate surface may change with temperature and residual gas incorporation, whilst Wilson and Weiss I6 have speculated that changes in film stress and/or silicon distribution may alter the distribution of planes with respect to the substrate surface and hence the ratio I,, ,/lzoo. which is measured parallel to the substrate surface.

Grovcnor er ~11%’ have discussed the variation of grain structure with temperature. summarising earlier work which developed

structure zone models based on the USC of a reduced temperature, r/T,,,. whcrc Tis the film tcmpcrature and r,,, is the melting point of the film. Modclling predicts that the variation of the grain diameter with time during deposition is initially linear then becomes quadratic as low energy grains cover the film surface’“. The effect of second phase particles, such as Si. AI,Cu or AI,Pd. can be to cause pinning of grain boundaries’” or. ifcertain criteria are met. abnormal grain growthJ”. The grain size may also be limited by the film thickness bccausc of grain boundary groove formation”.

Specular reflectivity

The specular reflectivity of Al and Al alloy films is often used as a measure of the film quality both to set up a deposition process and to monitor it. It is used because it is easy to assess cithcr visually or by a calibrated measurement. whilst it is also known

to affect subsequent photolithographic processing”‘.“. The specular reflectivity of a smooth aluminium sutfdce

between 100 and 1000 nm is shown in Figure 3. This has been plotted using the data of Hass and WaylonisJ” and Schultz”-‘” for the real and imaginary components of the refractive index. II and li, respectively, in Beer’s equation”

R = [(n-I)‘+k’]/[(n+ l)‘+k’]. (3)

Apart from an interband transition which reduces the specular reflectivity in the region of 800 nm. the specular reflectivity is large and fairly uniform over this wavelength range. The native oxide ofaluminium which is typically 40 8, thick. does not apprc- ciably affect the specular reflectivity above 200 nm”.

When the surface is no longer smooth. light is distributed away from the specular direction. Under conditions of near normal

721

R J W~/.sor~ and B L We/ss Alumlmum alloy fdm propertles

: i

24

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j I I I, l3W, - 126

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, I I I, (3lOJ . IX%1 , I I I, (700, -. :90 I I I I ) t2001 - -; I , I I I I (2001 . s:

, I I I, (100) . SY I I I I) (‘()I)) - I I I I I) (‘(JO) . IO I I I I) (200, . :o

RJ Wilson andB L Weiss: Aluminium alloy film properties

Absolute SDecular Reflectivitv % 95

94

93

92

91

90

89

88

87

86

d----Y

+ \ --+T + \i L+

85 200 300 400 500 600 700 800 900 1000

Wavelength In Nanometers

I Hass et al + Schultz

Figure 3. The specular reflectivity of a smooth Al surface.

The results show the importance of chamber base pressure and residual gases’.S’.“3. deposition temperature”.54. deposition pressure“.‘“, film thickness’.“.5J, alloy composition53.54, target to substrate distance” and number of passes under the target during deposition’. on the deposited films specular reflectivity. In addition to these factors the effect of substrate and substrate position in a batch deposition system have been shown to be important”. For sputtered films the change in specular reflectivity with deposition conditions has been explained using argu- ments relating to general surface roughness”,” or grain sizej’, usually on the basis of measurement at a single wavelength. However, Wilson and Weiss”’ have demonstrated that the specular reflectivity can depend on the surface roughness and surface optical properties only above a given wavelength, although below this wavelength it is also influenced by the film grain size which increases the specular reflectivity above that predicted from surface roughness considerations alone. It can be seen from equation (4) that this wavelength will depend on the relative contribution of the second term to the total reflectivity and hence on I-, 6 and A4. As discussed by Wilson and Weiss”, the specular reflectivity may be used as a deposition process parameter provided that it is recognised that it provides information which is directly relevant only to subsequent photolithographic processing.

The data suggest that there is a range of specular reflectivities which will give an optimum automatic alignment accuracy, and that the film specular reflectivity can be adjusted and kept within this range by suitable control of deposition parameters. The effect of specular reflectivity variation on critical dimension size42.4’ shows the importance of controlling the specular reflectivity uniformity in a batch deposition process. There is no proven causal connection between specular reflectivity and other film properties, and the tendency to link a high specular reflectivity with optimum film characteristics is not necessarily correct4’, although control of specular reflectivity uniformity is still impor-

tant, as is the identification of the causes of specular reflectivity change.

Resistivity

The resistivity of a deposited film used in silicon device tech- nology will affect the time delay introduced by the interconnect in the finished device. It is usually measured with a four point probe to determine the film sheet resistance, R,, and a surface profilometer to measure the film thickness, u’, the resistivity, p,, being given by

or = Ru’.

Chopras5 has comprehensively reviewed the literature relating to electron transport in metals. For a single crystal material the resistivity is determined by the amount of scattering caused by thermal vibrations of lattice atoms and by defects in the lattice. As the measurement temperature is reduced, the contribution to the resistivity caused by thermal vibrations decreases whilst the contribution due to defects remains unchanged. The temperature independent resistivity due to defects is termed the residual resistivity. In thin polycrystalline films there are two additional causes of electron scattering, namely grain boundaries and the film surface. The total resistivity of a thin polycrystalline film, p,, may be written as”

where p. is the resistivity within a grain (temperature dependent). B is the average grain diameter, d is the film thickness, L,, is the electron mean free path in the material (temperature dependent),

and y* and 7~ quantify the scattering process at grain boundaries and the surface, respectively. A resistivity ratio is often used to quantify the effects of deposition conditions on resistivity, and is defined as the ratio of the room temperature resistivity to that at a reduced temperature. typically that of liquid nitrogen (77 K) or helium (4.2 K), which gives increased sensitivity to residual

resistivity effects”.“. The resistivity of deposited Al alloy films has been seen to

depend on the partial pressure and type of residual gases present

during sputtering”.” “, the composition of the alloy”.‘4.5’~5x

and the deposition and annealing conditions’7.5x, the substrate bias voltage22, type of substrate for an Al-S1 alloy5’, substrate position in a batch deposition system” and the film thickness2’. The proposed physical mechanisms responsible for the changes are a possible change in grain boundary scattering with orientation’““, increased bulk conductivity due to gettering of

impurities at grain boundaries” and defects in the film caused by ion bombardment during bias sputtering” or electromigration”. The increase in resistivity associated with deposition either at poor base pressures or in the presence of controlled residual gas contamination is generally attributed to decreased film purity, without specifying whether it is caused by increased bulk resistivity or scattering at grain boundaries27,52.‘3. Verkerk et ~1” ” have shown that a large increase in film resistivity occurs when a granular structure develops for water partial pressures of 2 1.8 x 10m4 Pa, although the exact value is deposition rate

dependent, since this rate determines the relative amount of Al reaching the substrate compared with residual gaseous impurities. The incorporation of impurities was seen to be species dependent. with water tending to form oxide species at grain boundaries whilst oxygen tended to contaminate the whole film surface.

723

R J W//son and B L W.BSS Alumlnlum alloy film propertles

59 Al 1.5’1,,, SI ( IO(l) Si

The incrcascd rcsisti\ ity of’ Al lilms with alloying components

in the USC ofAlK%Ti films’” has been attributed to the deposited

alloy being a highly saturated solid solution in ;: mctartablc state

and hcncc containing cxccss lattice defects. Annealing produced

a Tao phase alloy cony-isina Al and AI,Ti prccipitatcs and

thercl‘orc a reduction in lattice dcfccts. A similar bchaviour has

been seen for Al Cu films”. whcrc ii Cu2AI phase is f-ormcd.

Wilson and Weiss<” have seen diffcrenccs in the sheet rcsistancc

01‘ Al l”/GSi films deposited onto (100) Si substrates Gth ot

without 21 I ~mi thick oxide layer deposited by 1.0~ Pressure

Chemical Vapour Deposition (LPC‘VD), attributing this to ;I

difkrcncc in the distribution of sihcon within the films. The

tcchniquc of resistance mcasurcmcnts has been applied to a study

of the precipitation of silicon in bulk “I’ and thin film Al Si”‘.

although the cfl‘ccts of silicon pr-ccipitatcd at grain boundal-1cs

724

on Frain boundary wattcring and grain go~‘t h ;irc ne~lcctccl.

The rcsi5ti\ity or sbcot rcsistancc 01‘ ;I dcpwitcd tilm may bc 01‘

use. as proposed by Fischer PI u/’ ‘. in predicting film rcliabilit!.

or in Idcntil’>ing tbc Icwl 01‘ precipitation ~ii Al Si alloys_ \IIICC‘

Si prccipitatc5 arc potential 5ites of failui-c”’ I”. In addition in ii

batch doposition system variation in lilm put-11) bc~wcn wh-

stratc positions may bc identified “I.

Film stress

‘The sukjcct of’tilm was and its measurcmcnt ha\ bwn rcvic,cd

by both Chopra” and HofTman”‘. The w-es\ can bc divided into

intrinsic and thcrnial cvniponcnts. Intrinsic 5trezs rrcstilts l‘rom

film strticlural considerations anti is ;I function 01‘ ticposition

conditions, \vhilc thermal stress is catwd b\, ;I mismatch in film

R J W//son and B L Wwss Alumlnlum alloy film properties

and substrate thermal expansion coctkients and occurs because

of changes in the sample tcmpcrature. If the film is unable to support the stress, then stress relief will occur. with the stress at which stress relief occurs, the stress relief mechanism and the position in the film of stress relief being dependent on the film structure. This subject is of relevance to silicon device manu- facture. since the mechanism of stress relief may lead to hillock formation which is detrimental to device reliability and, because

ofthe increasing importance of the failure mode. of stress induced void iormation.

For aluminium alloys on silicon or oxidised silicon substrates thermal stress dominates over intrinsic stress which may be neglected. This is becaube of the large diffcrcncc in thermal expansion coefficients bctwccn aluminium (23.5 x IO ’ K ’ at 20 C) and silicon (7.5 x IO ’ K ’ at 20 C) which results in the room temperature stress of an aluminium fihn deposited at a higher temperature being tcnsilc in the plane of the film. This causes the till71 -substrate combination to bow. as shown in Figure 4.

A number of investigations have been made of the variation of film stress with tcmpcrature” “. The results show remarkable similarity and a schematic stress- tcmpcrature diagram is shown in Figure 5. Point A is the tcnsilc stress in the film seen after deposition. On heating thcrc is a linear elastic reduction in tensile stress until, at temperature T,. the film is stress free. Further heating results in a linear elastic increase in compressive stress until point B is reached. Point B is the point at which the maximum compressive stress which can be supported by the film is reached and from B to C compressive stress relaxation occurs. On cooling from the maximum temperature at point C. there is an elastic reduction in compressive stress. with a stress free condition occurring at temperature T,. From T; to point D the tensile stress in the film incrcascs. until at point D. the maximuim tensile stress which can bc supported by the fihn is rcachcd and from D to E tensile stress relaxation occurs. The stress free tcmpcrature. T,. may be changed by stress relaxation at room tcmperaturc which rcduccs T, from its original value, which is visually assumed to have been the film temperature during deposition. or because of the deposition of a passivation layer which, ifundcr compressive stress, will induce an increased tcnsilc

stress in the fihn and hence incrcasc T,.

CompressIon Tension Stress

Distance

Figure 4. Schematic di:tgwm oE sample stress.

Typical values of the stress at room tcmpcraturc (A) and the temperature at which the stress is zero arc 100 200 MPa and 50& 100 C. respectively. The exact shape of the stress-tcmpcrature curve is dependent upon the stresses rcquircd to cause non-elastic bchaviour and the mechanism of stress relief”‘.““. The mechanism of stress relief has been the subject of speculation with grain boundary sliding”‘. grain boundary sliding and diffusional crcepfii, diffusional creep”“.” “. grain gro~,thhS.“‘.” and con- pound formation”’ being suggcstcd as the comprcssivc stress relief mechanism while dislocation slip”‘, dislocation loop gen- eration and grain boundary sliding” have been suggested as the tcnsilc stress relief mechanism. As discussed by Ashby”, thcrc are regions of tcmpcraturc and stress. which are dependent on film structure. in which different stress relief mechanisms arc likely to operate, whilst at boundaries ofthcse regions a combined mechanism may be operational. In addition suppression of one failure mode may lead to failure by another mode. and necessitate a different failure inhibition mechanism. Hillock growth is a rcsponsc to compressive stresses in the plant of the film. The non-uniform nature of hillock growth>’ tends to support the dual compressive stress relaxation mechanism of grain boundary sliding, which nucleates the hillock, and then difTusional creep “‘.“.” which transfers material to the hillock. allowing it to grow. The dif‘usional creep mechanism may be by matrix or grain boundary diffusion. although this is likely to bc tcmperaturc dependent.

The dcpendcncc of the stress rclicf mechanism on lilm struc- turchi ” ” ” highlights 21 need for control of film structure. if rcproduciblc bchavioul after tempcraturc cycling is to bc achieved. The addition of copper. which leads to the formation of AI?CLI precipitates at grain boundarics. rcduccs the tcndcncy of the film to form hillocks. this being due to the suppression of grain boundary sliding and/or suppression of grain boundary diffusion’“. Control of deposition conditions is still ncccssary to control the size and density of precipitates and their efI?ctivcncss in preventing hillock growth. Townsend and Vander PIas” have demonstrated the ability ofa Plasma Enhanced Chemical Vapou~ Deposited (PECVD) Si02 encapsulant. dcpositcd at 380 C. to suppress hillock growth in an AlL2”GCu film. and proposed that this was due to a change in the stress relief mechanism. From the stress tcmperaturc curve. it is sample tcmpcraturcs above T,

725

R J Wilsor? and B L We/ss Alumlmum alloy film propertles

R J Wilson and B L Weiss Alumlmum alloy film propertles

by deposited metal occurs. This leads to the conclusion that

step coverage can only be improved by mechanisms which cause rearrangement of material on the substrate such as ion bom-

bardment (bias sputtering), temperature or the removal of spccics which retard surface migration. Before considering the effects of substrate bias and temperature. some other factors which have been investigated will be discussed. Sputtering gas pressure”.” alloy composition, for alloys” containing up to I .2% Si and 4% Cu and the number of scans of the substrate past the target during deposition” have been shown not to affect step covcragc. It is also possible that the presence of residual gas species or surface impurities, due to perhaps poor cleaning techniques. could degrade step coverage.

Attempts at improving the step coverage of sputtered films by substrate heating. which do not involve the application of substrate bias, have been discussed by Inoue er crl” and Hart- sough and DenisonT3. Hartsough and Denison5’ controlled the film temperature by varying the deposition rate, using the heat of condensation of the arriving aluminium atoms. Improved step coverage was seen at increased temperatures. although only isolated steps were investigated. Inoue ct cd” investigated the step coverage of Al I ‘%Si alloy films deposited onto substrates with contact holes 0.9 /cm deep and of diamctcrs I.2 1.6 and 2.0 /mi etched into deposited oxide layers. The substrate was heated to its deposition temperature prior to deposition. lmprovcmcnt in the planarisation of the deposited film was seen as the tem- pcrature was increased from room temperature up to 565’ C. As the fihn becomes planar the step coverage can degrade. since the metal becomes thin at the top edge of the via. however. the deposition of more material results in a planar surface (Figure 6), and in such cases a tilling factor may be a more appropriate measure of the film profile than the step covcragc. In both cases” ” the improvement in the metallisation profile was attri- butcd to increased aluminium mobility at the substrate at elev- ated temperatures, A barrier layer is needed at high temperatures

to prevent reaction with the substrate. and mobility degradation. with TiN being shown to result in an improved more reproducible profile “.

The effect of applying a negative bias voltage to the substrate on the metal step coverage, has been discussed exten- sively2’.2’.“l r)j, Prior to reviewing the results. it should be remem-

d’

Typical Profile

step coverage = x/d

After Some Planar~satton The ProfIle Alters

step coverage = x’id

By Deposition Of More Material

step coverage = x”/d’

A Filling Factor d”/(d’+ z)

Mav NOW Be More Appropriate

bered that the effects of rf induced and dc negative bias voltages are not likely to be equivalent. especially if the substrate has an insulating layer present”‘, and that the energy of ions striking the substrate and their neutralisation energy will contribute to substrate heating”h. It is interesting to compare the results obtained by Smith et Al”‘.” and Park rt ~1“‘. When Smith et d” did not back-side coat the substrates prior to front-side mctallisation, the results showed considerable variation compared with the back-side coated case. Both Homma and Tsunckawa”’ and Hoffman et d’” have investigated films sputtered using rf induced negative substrate bias. Homma and Tsunekawa”’ showed that considerable resputtering of material (70%) was required to cause an aluminium film to planarise in 1.5 ilrn wide and deep grooves. HofIman et ~1”’ applied rf bias to substrates held at 480 C and also saw planarisation over 2.0 ktrn wide. 1 {trn deep grooves. They noted that contacts proved harder to fill than grooves and that features larger than the ones planarised were not fully planarised. A TiSiz 4 layer was required to prevent reaction with the substrate. The proposed mechanisms for the planarisation were resputtering”’ and increased diffu- sivity”“. However, Homma and Tsunckawa” did not eliminate temperature as a factor. since their substrates achieved temperatures up to 250 C. whilst Hoffman ct crl”“did not separate the effects of rf bias and temperature and did not report resputtering.

No rcsputtcring has been reported to occur during the application of negative dc bias to back-side coated substrates” ‘4.“‘. Skclly and Gruenke”’ showed that the bias voltage was changing the energy of ions incident on the substrate at voltages beyond -20 V. Problems arose in attempting to planarise a range 01 fcaturc sizes. Most of the via filling took place after an induction period, which suggested a tcmpcraturc threshold. although cfficts of temperature and ion bombardment induced mobility were not scparatcd. Some fotward sputtering of material alas noticed during the early stages of deposition. Part c’/ crl” compared the cffccts of two target designs which rcsultcd in 0.53 W cm ’ (FOCEST) and 0.15 W CIII ’ (INSET) bias power density at - 150 V dc bias. The higher power density provided planarisation” compared with a slight step covcragc increase for the low power density” 14. The work of Smith” and Park ct da4 shows that temperature or ion bombardment on their own arc not suficient to cause improved step covcragc;planarisation in the region of deposition conditions investigated. A mechanism involving thermal activation of defects induced by ion bombardment was proposed.

Other techniques such as LPCVD may bc able to provide conformal coverage of features but suffer from other dis- advantages rclatcd to the deposition method”‘. In addition, a via plug of the same material as used for mctallisation may not provide an optimum redistribution of current through the via”“. To avoid current crowding a via or contact should be filled with a high resistivity plug, such as tungsten.

Conclusion

This paper has rcvicwcd the properties of deposited Al alloy films. The properties chosen were electromigration rcsistancc, specular reflectivity. resistivity, stress and step coverage because these constitute areas of concern during device processing. An attempt has been made to illustrate links between the film structure obtained under a given set of deposition conditions and the film properties. while bearing in mind the rclcvance of the information to deposition monitoring.

727

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RJ Wilson and 6 L Weiss: Aluminlum alloy film propertles

” C Li. R D Black and W R La Fontaine, Appl Ph~s Lef/. 53,31 (1988). “Y Koubuchi. J Onuki. M Suwa and S Fukada, Pruc 6th In/ IEEE VLSI Multilrwl Interconnwtion Cor$ p 419. IEEE, New York (1989). *‘J G Ryan. J B Riendeau, S E Shore, G J Slussen, D C Beyar, D P Bouldin and T D Sullivan, J Vuc Sci Techno/, A8, 1474 (1990). X4 W Tice and G Slusser, J Vuc Sci Technol, BS, 106 (1990). “K Hinode, I Asano, T lshiba and Y Homma, J Vuc Sci Technol, B8, 495 (1990). “M Kato, H Niwa, H Tagi and H Tsuchikawa, J Appl Ph~‘.s. 68, 334 (1990). “H Bangert. A Kaminitscheik, A Wagendristel, A Barna. P B Barna and G Radnoczi, Thin Solid Films. 137, 193 (1986). xX B Jonsson and S Hagmark, Thin Solid Films. 114, 257 (1984). x9 D Barton and C Maze, Srmicond In/. 8, 98 (I 985).

““C J Dell’oca, Proc 2nd ht IEEE VLSI Mulrikcel Inierconnrcrion Con/; p 3. IEEE, New York (1985). ” Y Homma and S Tsunekawa, .I Ekfrochem Sot, 132, I466 ( 1985). ‘ID W Skelly and L A Gruenke, J Vu .Sc,i Techno/, A4,457 (1986). “J F Smith, F T Zold and W Class, Thin Solid Films. 96,291 (1982). ““V Hoffman, J Griswold, D Mintz and D Harra, Thin Solid Films, 153, 369 (1987). “‘J J Cuomo, R J Gambino and R Rosenberg, J Vu SCI Twhnol. I I, 34 (1974). “A N Pargellis, J Vu SC; Twhnol, A7, 27 (1989). “R A Levy, M L Green and P K Gallagher. J Elwrrochcw Sot, 131, 2175 (1984). ‘lx J M Pimbley and D M Brown. lEEE Trtrn.v Ekw/ron Dwiw.s, 33, I399 (1986).

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Documents

A review of the properties of aluminium alloy films used during silicon device fabrication