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Dense fully 111-textured TiN diffusion barriers: Enhanced lifetime throughmicrostructure control during layer growthJ.-S. Chun, I. Petrov, and J. E. GreeneCitation: J. Appl. Phys. 86, 3633 (1999); doi: 10.1063/1.371271View online: http://dx.doi.org/10.1063/1.371271View Table of Contents: http://jap.aip.org/resource/1/JAPIAU/v86/i7Published by theAmerican Institute of Physics.Related ArticlesEvaluation and modeling of lanthanum diffusion in TiN/La2O3/HfSiON/SiO2/Si high-k stacks

Appl. Phys. Lett. 101, 182901 (2012)Influence of catalytic gold and silver metal nanoparticles on structural, optical, and vibrational properties of silicon

nanowires synthesized by metal-assisted chemical etchingJ. Appl. Phys. 112, 073509 (2012)Chemical and structural investigations of the interactions of Cu with MnSiO3 diffusion barrier layersJ. Appl. Phys. 112, 064507 (2012)Thermal tweezers for nano-manipulation and trapping of interacting atoms or nanoparticles on crystallinesurfacesJ. Chem. Phys. 137, 114701 (2012)Thermal conductance at atomically clean and disordered silicon/aluminum interfaces: A molecular dynamicssimulation studyJ. Appl. Phys. 112, 054305 (2012)Additional information on J. Appl. Phys.Journal Homepage: http://jap.aip.org/Journal Information: http://jap.aip.org/about/about_the_journalTop downloads: http://jap.aip.org/features/most_downloadedInformation for Authors: http://jap.aip.org/authors

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Dense fully 111-textured TiN diffusion barriers: Enhanced lifetime throughmicrostructure control during layer growth

J.-S. Chun, I. Petrov, and J. E. Greenea)

Materials Science Department, Materials Research Laboratory, and Coordinated Science Laboratory,University of Illinois, 1101 West Springfield Avenue, Urbana, Illinois 61801

Received 16 April 1999; accepted for publication 24 June 1999

Low-temperature deposition of TiN by reactive evaporation or sputter deposition onto amorphoussubstrates leads to highly underdense layers which develop mixed 111/002 orientations through

competitive growth. In contrast, we demonstrate here the growth of low-temperature 450 C fully

dense polycrystalline TiN layers with complete 111 texture. This was achieved by reactive

magnetron sputter deposition using a combination of: 1 highly oriented 25-nm-thick 0002 Ti

underlayers to provide 111 TiN orientation through texture inheritance local epitaxy and 2 high

flux (JN2/JTi14), low-energy (EN

220eV), N2

ion irradiation in a magnetically unbalanced

mode to provide enhanced adatom diffusion leading to densification during TiN deposition. The Ti

underlayers were also grown in a magnetically unbalanced mode, in this case with an incident

Ar/Ti flux ratio of 2 and EAr11 eV. All TiN films were slightly overstoichiometric with a N/Ti

ratio of 1.020.03. In order to assess the diffusion-barrier properties of dense 111-textured TiN, Al

overlayers were deposited without breaking vacuum at 100 C. Al/TiN bilayers were then annealed

at a constant ramp rate of 3 C s1 to 650 C s1 and the interfacial reaction between Al and TiN

was monitored by in situ synchrotron x-ray diffraction measurements. As a reference point, we findthat interfacial Al3Ti formation is observed at 450 C in Al/TiN bilayers in which the TiN layer is

deposited directly on SiO2 in a conventional magnetically balanced mode and, hence, is underdense

with a mixed 111/002 orientation. However, the onset temperature for interfacial reaction was

increased to 610 C in bilayers with fully dense TiN exhibiting complete 111 preferred orientation.

1999 American Institute of Physics. S0021-8979 99 04119-5

I. INTRODUCTION

Al-based interconnects with TiN diffusion barriers are

widely used in multilevel metallization architectures incor-

porated in submicron ultralarge-scale integrated ULSI de-

vices. As the dimensions of integrated circuits continue to

decrease while their packing density increases, interconnec-

tion reliability issues such as electromigration resistance and

diffusion barrier lifetime become increasingly important. It is

well known1 5 that electromigration lifetimes are enhanced

in Al metal lines having a large average grain size, a small

grain-size distribution, and strong 111 texture. Shibata et al.6

found that the dominant factor for Al linewidths below 0.5

m is the degree of 111 preferred orientation.

Texture control can also be important for controlling the

stability of heterointerfaces during high temperature process-

ing. Both the rate and extent of reaction at metal/diffusionbarrier interfaces have been shown to depend upon film

texture.7,8 In all previous studies718 of polycrystalline Al/

TiN interfacial reactions, there were no attempts made to

control TiN film texture or, except for Refs. 7 and 9, poros-

ity. In fact, the TiN layers are often uncharacterized. Under

typical industrial deposition conditions, sputter-deposited

TiN is underdense with a mixed 111/002 texture.4,7 Al/TiN

interfacial reactions, the formation of Al3Ti, and barrier fail-

ure occur during isothermal annealing for 30 min at tempera-

tures below 400 C.19 Air-exposure leading to oxide forma-

tion prior to Al deposition has been shown to increase the

thermal stability of the barrier layer.19 However, the resultsare not reproducible. Moreover, air-exposure reduces the de-

gree of Al 111 texture and increases the contact resistance.

Greene et al.20,21 has shown that the preferred orienta-

tion of polycrystalline TiN films grown by ultrahigh-vacuum

UHV reactive-magnetron sputter deposition on amorphous

SiO2 at 350C in pure N2 discharges can be controllably

varied from essentially complete 111 to purely 002 by vary-

ing the incident ion/metal flux ratio Ji/JTi from 1 to 5

while the ion energy EN2 is maintained constant at 20 eV.

The incident ions are predominantly N2 96.3% ,22 corre-

sponding to 10 eV per incident N, well below the lattice

displacement threshold. During low-temperature deposition

in the absence of significant ion irradiation (Ji/JTi1), TiN

initially exhibits a mixed texturepredominantly 111, 002,

and 022and then slowly evolves to a strong 111 texture

containing a network of both inter- and intracolumnar poros-

ity due to the combination of low adatom mobilities and

competitive columnar growth. Increasing Ji/JTi5 results in

enhanced adatom mobilities giving rise to a dense micro-

structure. In addition, the low-surface-energy 002 texture

now dominates even at submonolayer coverage and there isa Electronic mail: [email protected]

JOURNAL OF APPLIED PHYSICS VOLUME 86, NUMBER 7 1 OCTOBER 1999

36330021-8979/99/86(7)/3633/9/$15.00 1999 American Institute of Physics


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no longer any indication of competitive growth as 002 col-

umns grow essentially straight up.20,21

Obtaining a dense TiN layer with complete 111 texture

at low growth temperatures, as required for diffusion barrier

and many hard-coating applications, is a challenging growth

kinetics problem. Moving toward higher growth tempera-

tures favors the development of the 002 texture,23 rather than

111, since 002 is the low-energy surface for TiN.24 In ad-

dition, the use of ion irradiation to promote densification of111-textured TiN leads to either 002 preferred orientation as

noted above for low ion energies with high ion-to-neutral

ratios20,21 or very high in-plane compressive stress and

mixed texture with higher energies and low ion-to-neutral

ratios.23

In this article, we show, for the first time, that fully

dense completely 111-textured TiN layers can be achieved at

low growth temperatures using a combination of 1 highly

oriented 25-nm-thick 0002 Ti underlayers to provide 111

orientation through texture inheritance local epitaxy and 2

high-flux (JN2/JTi14), low-energy (EN

220eV), N2

ion

irradiation to provide enhanced adatom diffusion leading to

densification. Both the Ti and TiN layers were grown in a

magnetically unbalanced mode. While interfacial Al3Ti for-

mation was observed at 450 C in Al/TiN bilayers with un-

derdense TiN having a mixed texture, the onset temperature

was increased to 610 C for bilayers with fully dense 111

TiN. The Al overlayers are dense in both cases, but exhibit

much stronger 111 texture i.e., texture inheritance on dense

111 TiN.

II. EXPERIMENTAL PROCEDURE

All films were grown in a load-locked multichamber

UHV stainless-steel dc magnetron sputter deposition systemwhich has been described in detail elsewhere.25 The pressure

in the sample introduction chamber was reduced to less than

5108 Torr 7106 Pa) prior to initiating substrate ex-

change into the deposition chamber which has a base pres-

sure of 51010 Torr 7108 Pa). For the present experi-

ments, an additional magnetron-sputtering source with

separate water cooling lines and shutter was added in the

viewport flange opposing the original source. The Ti target

99.999% pure was operated in a magnetically unbalanced

mode for the growth of both Ti and TiN films while the Al

target 99.999% pure was sputtered in a magnetically bal-

anced mode. Target-to-substrate separations were 6.5 cm for

TiN and 10 cm for Al deposition.A pair of external Helmholtz coils with Fe pole pieces

were utilized to create a uniform axial magnetic field Bextin the region between the target and the substrate. The posi-

tive and negative signs refer to geometries in which B ext aids

and opposes, respectively, the field of the outer pole of the

magnetron. B ext has a strong effect on the ion flux incident at

the substrate, with only minor effects on the target atom

flux.25

The substrates used in these experiments were thermally

oxidized 11 cm2 Si 001 wafers with an SiO2 thickness of

0.6 m. They were cleaned with successive rinses in ultra-

sonic baths of trichloroethane, acetone, ethanol, and deion-

ized water and blown dry with dry N2 . The wafers were

mounted on resistively heated Ta platens using Mo clips and

inserted into the sample introduction chamber for transport

to the growth chamber where they were thermally degassed

at 800 C for 1 h. The substrate temperature Ts was then

adjusted to 450 C as the Ti target was sputter etched for 5

min while shielding the substrate and Al target. Reported Tsvalues, which were calibrated using a 0.25-mm-diam

chromelalumel thermocouple bonded to the surface of asacrificial substrate using Ag paste, include the contribution

from plasma heating.

Plasma characteristics in the vicinity of the substrate

were determined as a function of B ext from electrical mea-

surements, following the procedures described in Refs. 25

and 26, obtained using both disk-shaped and cylindrical

probes. Most of the ions incident at the substrate experience

the full substrate sheath potential Vs(VbVp), where Vpis the plasma potential and Vb is substrate bias, since the

mean-free path for charge exchange collisions,27 8 mm, is

more than an order of magnitude larger than the sheath

width,

28

estimated from the ChildLangmuir equation. In thepresent experiments, Vb is equal to the floating potential Vfand the ion energy Ei at the substrate is e ( VfVp ).

Two sets of TiN layers were grown with Ts450 C at

the relatively high sputtering pressure of 20 mTorr

N2(99.999%), chosen to suppress kinetic energy transfer to

the growing film by particles backscattered from the target.

In sample set i, TiN layers, 160 nm thick, were deposited on

SiO2 with JN2/JTi1 and EN

220eV (Vp3 V and Vf

17V). The ion current density at the substrate,

j N20.12 mA cm2 corresponding to an ion flux

JN27.51014 cm2 s1), was obtained using an external

axial magnetic field B ext of20 G. These conditions resultedin a deposition rate R of 6.0 nm min1. The second set of

TiN layers, series ii, were grown on 25-nm-thick Ti under-

layers deposited on SiO2 at Ts80 C in 20 mTorr Ar

99.999% with an incident Ar/Ti flux ratio of 2

(JAr6.11015 cm2 s1 obtained with Bext200 G), and

an Ar ion energy EAr11 eV (Vp12.5 V and Vf23.5 V). The TiN overlayers, 140 nm thick, were then

deposited without breaking vacuum using JN2/JTi14

(JN21.61016 cm2 s1 achieved with B ext150 G) and

EN223eV (Vp11V and Vf34 V). The deposition

rates R of Ti and TiN were 30 and 5.4 nm min1, respec-

tively.Calculations carried out using the TRIM90 computer

code,29 combined with Monte Carlo gas transport simula-

tions based upon energy-dependent scattering cross

sections,30 show that for the TiN growth conditions de-

scribed above, a majority of both the sputter-ejected species

and the energetic neutrals backscattered from the target are

thermalized during gas-phase transport and contribute little

to the total kinetic energy transfer to the substrate and grow-

ing film. The ion flux to the substrate is due, as determined

by double-modulated mass spectroscopy measurements,22

primarily to N2(96.3%) with the remainder being

N(3.5%) and Ti(0.2%).

3634 J. Appl. Phys., Vol. 86, No. 7, 1 October 1999 Chun, Petrov, and Greene


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Immediately following TiN deposition, for both series i

and ii, Ts was decreased to 100 C and 160-nm-thick Al

overlayers were grown in 5 mTorr Ar discharges without

breaking vacuum. The Al deposition rate was 71.1 nm min1

with an incident Ar/Al flux ratio of 0.8 and EAr10 eV.

The microstructure and microchemistry of as-deposited

and annealed samples were determined using a combination

of Rutherford backscattering spectroscopy RBS , x-ray dif-

fraction XRD , and plan-view and cross-sectional transmis-sion electron microscopy TEM and XTEM . The RBS probe

beam consisted of 2 MeV He ions incident at an angle of

22.5 relative to the sample surface normal with the detector

set at a 150 scattering angle. Backscattered spectra were

analyzed using the RUMP simulation program.31 The uncer-

tainty in determining the N/Ti ratio was less than 0.03. TiN

exhibits a wide single-phase field with a N/Ti ratio extending

from0.6 to 1.2.32 In the present experiments, all TiN films,

irrespective of Ji/JTi , were found to be slightly oversto-

ichiometric with N/Ti1.020.03, as expected for low-

temperature growth in pure N2 discharges.

XRD 2 and -rocking curves were obtained using

Cu K radiation in a system equipped with a double-crystalspectrometer to provide a resolution of 0.01 2 in the pow-

der diffraction mode 45 kV, 20 mA, 1 divergent slit . The

TEM and XTEM analyses were carried out using Philips

CM12 and Hitachi 9000 electron microscopes operated at

120 and 300 kV, respectively. Sample preparation for TEM/

XTEM examination followed the procedure outlined in Ref.

21.

The surface morphologies of as-deposited TiN films

were investigated using atomic force microscopy AFM .

The measurements were carried out in a Digital Nanoscope

II microscope operated in air in the contact mode with oxide-

sharpened Si3

N4

tips having radii of 540 nm.

In situ x-ray diffraction XRD measurements during

sample annealing were performed at the Brookhaven Na-

tional Synchrotron Light Source, beamline X-20C. As-

deposited samples were loaded in an annealing chamber

aligned to the beamline and equipped with an x-ray transpar-

ent Be window. The chamber was evacuated to 5

106 Torr, backfilled with purified He, re-evacuated, and

then backfilled with 1 atm of He. As-deposited samples were

annealed using linear temperature ramps from 100 to 650 C

at 3 C s1. A high-intensity monochromator employing

bandpass multilayer filters provided an energy resolution of

1.5% at 6.9 keV 0.1797 nm with a typical intensity at the

sample of 3

10

12

photons s

1

. The incident x rays illumi-nated a sample area 15 mm2. Diffraction scans were col-

lected using a position-sensitive detector with time resolution

of ms.

III. EXPERIMENTAL RESULTS AND DISCUSSION

Figure 1 a is an XTEM micrograph from a TiN layer

grown directly on SiO2 with JN2/JTi1 and EN

220 eV,

conditions corresponding to TiN sample series i. The micro-

graph is typical of TiN films grown under weak ion irradia-

tion conditions at Ts500C (Ts /Tm0.2). The layer is

underdense due to limited adatom mobilities leading to

atomic shadowing which, in turn, results in a columnar mi-

crostructure with both inter- and intracolumnar voids.20,21

Careful analyses, using nanodiffraction and high-

resolution XTEM, of the region near the film/substrate inter-

face show that the layer initially nucleates with islands lead-

ing to grains which are predominantly broad-based 002 and

narrow 111. However, the 111 grains eventually overgrow

the 002 via a competitive growth mechanism. That is, diffu-sion on the low-energy 002 surface24in which there is only

one adatom backbondis fast, leading to islands spreading

out relatively rapidly and accounting for the broad-based 002

columns, compared to the 111 surface where each adatom

has three backbonds. Thus, the steady-state adatom density

on 002 islands is low as adatoms quickly move to the edges,

while 111 grains grow upwards rapidly and intercept ever

larger fractions of the incoming flux. While the majority of

the 002 grains are overgrown early in the growth process,

some survive even in films with thicknesses up to 1 m.21

The open intercolumnar boundaries observed in low-

temperature 111-oriented transition-metal nitride layers e.g.,

TiN,20,21 Ti1xAlxN, 33 and ScN34 is simply a consequenceof the effect of the atomic shadowing process during texture

evolution via competitive growth combined with the distri-

bution of particle incidence angles. Figure 1 a also shows

that the sample surface is rough and faceted, as commonly

observed for films with voided column boundaries. The layer

thickness determined from the micrograph is 160 nm, in

good agreement with the deposition rate calibration.

Figure 1 b is a -2 XRD scan from the series i layer

shown in Fig. 1 a . The pattern contains an intense 111 peak

and a small 002 peak, both at positions corresponding to

B1-NaCl-structure TiN. The integrated 111 to 002 peak in-

tensity ratio I111/I002 is 8.8, a factor of 14 higher than that

FIG. 1. a Bright-field XTEM micrograph and b corresponding -2XRD

scan from an as-deposited series-i TiN layer grown on SiO 2 at 450 C with

JN2 /JTi1 and EN

220 eV.

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obtained from randomly oriented TiN powder. Thus, the

dominant texture at this film thickness is 111 in agreement

with the XTEM results discussed above. The 111 interplanar

spacing d111 is 0.244 nm which is close to the reported bulk

value of 0.2449 nm and indicates that the layer is nearly fully

relaxed. The full width at half maximum FWHM intensity

of the 111 -rocking curve, not shown, is 23.1. The

large value of is consistent with kinetically limited com-

petitive texture evolution.

Low-energy, high-flux, N2 irradiation during growth

densifies the TiN microstructure leaving no residual ion-

induced defects detectable by XRD and TEM.20,21 However,

the layers also exhibit complete 002 preferred orientation.Thus, in an attempt to obtain TiN layers which are both

dense and have complete 111 texture, we have carried out

experiments combining intense N2 ion irradiation with the

use of a dense 0002 oriented hcp Ti underlayer as a crystal-

lographic template sample series ii . The hcp Ti basal plane

and the fcc TiN 111 plane the NaCl crystal structure is fcc

with a two-atom basis have the same two-dimensional sym-

metry. Moreover, Ti underlayers are often used in multilevel

interconnections to provide enhanced adhesion between the

diffusion barrier and SiO235 and lower contact resistance

with Si.36

Figure 2 is a schematic view of the atomic arrangement

at the interface between 0002 Ti and 111 TiN layers. The

misfit along equivalent close-packed directions in TiN 111

and Ti 0002 is approximately 1.6% with xTiN0.29995 nm

and x Ti0.29503 nm. 111 is a polar direction in TiN with

alternate 111 planes consisting of all Ti and all N atoms.

Preferred orientation in metal films, as opposed to

transition-metal nitride layers, typically corresponds to the

low-energy surface even at low growth temperatures37 due to

FIG. 2. Schematic view of the atomic arrangement at the interface between

Ti 0002 and TiN 111 layers.

FIG. 3. a -2 XRD scan of an as-deposited Ti layer grown on SiO2 at80 C with an incident Ar/Ti flux ratio of 2 and EAr11 eV. b An 0002

-rocking curve and c a bright-field plan-view TEM micrograph from the

same sample.



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much higher adatom mobilities on metal surfaces. In thepresent experiments, Ti underlayers grown on SiO2 at 80 C

(Ts /Tm18%) were found to be 0002. However, to enhance

the degree of orientation, we employed low-energy Ar ir-

radiation through the use of magnetically unbalanced sputter

deposition to provide, at the growth surface, an ion-to-metal

flux ratio JAr/JTi2 with EAr11 eV.

Figure 3 a is a -2 XRD scan of an as-deposited Ti

layer. The only peaks observed are 000l ( l2,4,...) whereas

for Ti powder patterns the strongest peak is 1011 2

40.17 with an intensity ratio I1011 /I00023.4. We obtain

an 0002 interplanar spacing d0002 of 0.2342 nm in agreement

with the reported bulk value, 0.2341.38 The -rocking curve

shown in Fig. 3 b has a FWHM value 3.3. A bright-

field plan-view TEM micrograph of the same sample isshown in Fig. 3 c . The absence of Moire fringes is indica-

tive of a columnar structure. However, the individual grains

appear fully dense with no evidence of either inter- or intra-

grain porosity. The average Ti grain size obtained from plan-

view TEM images using image analysis software39 is 2621

nm.

Immediately following growth of the oriented Ti under-

layers, TiN was deposited at P N220mTorr and Ts

450 C with JN2/JTi14 and EN

220 eV. Figure 4 a is

a typical bright-field XTEM micrograph from an as-

deposited TiN/Ti bilayer structure. TiN and Ti layer thick-

nesses determined from the micrographs are 160 and 5 nm,

FIG. 4. a Bright-field XTEM micrograph from an as-deposited series-ii

TiN/Ti bilayer structure. The TiN layer was grown at 450 C with

JN2 /JTi14 and EN

220 eV. b A -2 XRD scan and c a 111

-rocking curve from the same sample.

FIG. 5. High-resolution XTEM micrograph of the TiN/Ti interface from the

sample corresponding to Fig. 4.



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respectively. The decrease in the Ti thickness from 25 to 5

nm indicates that 80% of the underlayer has reacted with

the incident N2 flux, due to the high N2/Ti ratio used in these

experiments, to form TiN during the initial stages of nitride

deposition. There are two important points to note in this

micrograph. The first is that this TiN layer, in contrast to the

sample in Fig. 1 a , is dense with no detectable inter- or

intracolumnar porosity. The film appears fully dense even in

high-resolution micrographs. The second point is that Fig.

4 a reveals no indication of competitive 111/002 grain

growth. In fact, all TiN columns are 111 and grow commen-

surately with the underlying transformed TiN grains as will

be shown below using high-resolution XTEM HR-XTEM .

Figure 4 b is a -2XRD scan from the type-ii bilayer

sample corresponding to Fig. 4 a . Consistent with the

XTEM results, and in contrast to XRD scans from the type-i

TiN layer grown with no intentional ion irradiation and with-

out the oriented Ti underlayer, only the 111 peak is obtainedwith no evidence of a residual 002 feature. The d111 lattice

spacing is 0.2442 nm in reasonable agreement with the bulk

value indicating that the layer is relaxed. XRD pole figures

not shown obtained with 2 set to the 111 Bragg angle for

TiN confirm that the film texture is azimuthally symmetric

i.e., it exhibits fiber texture . No other diffraction peaks

were detected in either pole figures or -2 scans. The lack

of Ti reflections is due to the fact that most of the underlayer,

as shown in the XTEM micrograph in Fig. 4 a , was trans-

formed to TiN. The 111 -rocking curve in Fig. 4 c is ex-

tremely narrow, 1.9, compared to the FWHM obtained

from the type-i layer, 23.1. Thus, the mosaicity is

greatly reduced giving rise to a factor of16 increase in thein-plane coherence length.

A HR-XTEM bright-field image of the TiN/Ti interface

in Fig. 4 a is presented in Fig. 5. It shows clearly, together

with high and lower-magnification images obtained from dif-

ferent regions of the same sample, that cubic 111 TiN grains

grow with a local epitaxial relationship to hexagonal 0002 Ti

grains: TiN111Ti0002 . The TiN/Ti interface is locally coher-

ent with TiN and Ti exhibiting equal in-plane grain sizes.

The micrograph also shows that the Ti/SiO2 interface re-

mains abrupt.

Figures 6 a and 6 b are plan-view TEM micrographs

showing a direct comparison of the microstructure of type-i

FIG. 6. Bright-field plan-view TEM micrographs of a series-i and b

series-ii TiN/Ti bilayers.

FIG. 7. AFM images of the samples shown in a Fig. 6 a and b Fig. 6 b .



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TiN films grown by conventional reactive magnetron sput-

tering and type-ii TiN layers grown under intense low-

energy N2 irradiation and with an oriented Ti underlayer.

Both TiN films were deposited at Ts450 C with

EN220 eV. However, the series-i film in Fig. 6 a grown

with JN2/JTi1 exhibits an underdense structure consistent

with the XTEM micrograph in Fig. 1 a with a mixed tex-

ture which is predominantly 111. The intercolumn voids

have widths of up to 2 nm. In contrast, the series-ii layer in

Fig. 6 b is fully dense with a complete 111 texture. Averagegrain sizes d measured near the top of the underdense anddense TiN films are 2014 and 4330 nm, respectively.

Since the average grain size of the as-deposited Ti under-

layer is 2621 nm, d in the 111 TiN overlayer grainsincreases with layer thickness. This can be also seen in the

XTEM micrograph in Fig. 4 a .

Figures 7 a and 7 b are AFM images of the TiN films

corresponding to Figs. 6 a and 6 b . The 160-nm-thick un-

derdense series-i TiN layer in Fig. 7 a is extremely rough

with a root-mean-square surface width w 8.5 nm and anin-plane surface coherence length 78 nm. The same

thickness dense series-ii 111-oriented TiN layer Fig. 7 b

grown on the 0002 Ti template is much smoother with w

3.3 nm and 49 nm.In order to examine the effects of texture, microstruc-

ture, and surface roughness on the diffusion barrier proper-

ties of TiN, 160-nm-thick Al overlayers grown with

JAr/JAl0.8 and EAr10 eV were deposited at

PAr5 mTorr and Ts100 C on both series-i and series-ii

TiN layers without breaking vacuum. Figures 8 a and 8 b

are typical XTEM micrographs from Al/TiN diffusion-

barrier test structures. Both Al overlayers were fully dense

with average grain sizes, obtained from plan-view images, of

210 nm. They exhibited quite different degrees of pre-

ferred orientation, however, and the surface of the Al film

grown on the type-i TiN layer is much rougher with an in-

FIG. 8. Bright-field XTEM micrographs from a series-i Al/TiN and b

series-ii Al/TiN bilayers.

FIG. 9. -2 XRD scans from as-deposited a series i and b series-ii

Al/TiN bilayers. c Al 111 -rocking curve from the sample corresponding

to b .



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plane wavelength of the order of the grain size.

-2 XRD scans from the Al/TiN bilayers in Figs. 8 a

and 8 b are shown in Figs. 9 a and 9 b . The patterns dem-

onstrate that the Al film grown on TiN with mixed 111/002

texture also exhibits a mixed orientation Fig. 9 a while the

Al layer grown on the TiN film with a complete 111 pre-

ferred orientation also has a pure 111 texture note the scale

change by a factor of 102 in the inset of Fig. 9 b . Both Al

layers are relaxed with a 111 interplanar spacing equal to thebulk value of 0.2338 nm.40 In addition, the 111 Al and TiN

integrated peak intensities obtained from bilayer ii are two

orders of magnitude higher than those of bilayer i at equal

layer thicknesses. for Al layers grown on both series i

and ii TiN are 13.5 and 0.82, respectively.

The above results show that the stronger 111 orientation

in the TiN film in bilayer-ii results in a corresponding en-

hancement in the 111 texture of the Al overlayer. This occurs

through texture inheritance as metal islands nucleate and, at

least initially, grow with a local epitaxial relationship with

underlying TiN grains. However, the final Al grain size is

more than five times larger than that of TiN indicating the

presence of recrystallization which is known to occur at tem-peratures Tr as low as Tr /Tm0.24 in high purity Al.

41

While this further reduces grain misorientation in both Al

overlayers, the mosaicity of the layers grown on purely 111-

oriented TiN is still more than an order of magnitude lower,

as indicated by the measurements, than for Al grown on

mixed 111/002-oriented TiN.

Figures 10 a and 10 b are typical contour plots of the

logarithm of the diffracted synchrotron x-ray intensity as a

function of the diffraction angle 2 and temperature Ta dur-

ing thermal ramping of underdense series-i TiN and dense

111 series-ii TiN samples to 650 C at a rate of 3 C s1. The

intensity is indicated by gray scale. White corresponds to an

intensity of 1.9104 in Fig. 10 a and 9.2104 in Fig. 10 b

while black is 102 counts s1. The angular range of the

position sensitive detector was chosen based upon ex situ

XRD results following isothermal anneals showing, in addi-

tion to the strong TiN and Al 111 peaks, the emergence of

the 112 tetragonal-structure Al3Ti peak also observed by

TEM and XTEM at 39.132. 112 was the only XRD peak

observed, indicating that the product phase is highly tex-

tured. Grains of the intermetallic compound Al3Ti, whichhas the DO22 structure with lattice constants a0b 00.385 37 nm and c0.858 39 nm,42 are aligned with their

112 plane perpendicular to the Al film growth direction.

112 is the closest-packed Al3Ti plane and exhibits an in-

plane misfit with Al 111 of only 1%.

Figure 10 a shows that for type-i samples, the intensity

of the 111 Al peak at 45.2 0.1797 nm 2decreases with

increasing annealing temperature Ta and is completely lost at

Ta620 C. This is well below the Al melting point

660 C and indicates that the Al layer has been completely

consumed. Al3Ti starts to form at 450 C and reaches its

maximum intensity at the same temperature at which the Al

intensity disappears. In contrast, Al3Ti formation in type-iisamples is suppressed up to 610 C and significant unre-

acted Al still remains even at 650 C as shown in Fig. 10 b .

We attribute the enhanced interfacial stability of series

ii, compared to series i, Al/TiN samples to the densification

of TiN grain boundaries with high-flux (JN2/JTi14), low-

energy (EN220eV), N2

ion irradiation during TiN depo-

sition, the larger TiN grain size by a factor of two , and the

flatter surface. It has been reported previously8,9 that TiN

barrier failure is initiated at the Al/TiN interface by the for-

mation of Al3Ti in the Al layer. In type-i samples, massive

diffusion of Al into the open TiN inter- and intracolumnar

voids is expected to occur rapidly at relatively low annealingtemperatures leading to an enormous increase in the actual

area of the Al/TiN interface.

IV. CONCLUSION

We have demonstrated, for the first time, the low-

temperature growth of fully dense polycrystalline TiN with

complete 111 preferred orientation. This was achieved using

a combination of highly oriented thin 0002 Ti underlayers to

provide orientation through texture inheritance local epi-

taxy and high-flux, low-energy, N2

ion irradiation(JN

2/JTi14, EN

2 20 eV) to provide enhanced adatom

diffusion leading to densification. The preferred orientation

of Al overlayers grown on dense fully 111-oriented TiN was

also greatly enhanced due to texture inheritance, thus reduc-

ing the FWHM of the 111 -rocking curve from 13.5 for Al layers grown on TiN deposited by conven-

tional reactive magnetron sputter deposition to 0.82.

This corresponds to an enhancement in the average mosaic

coherence length by more than a factor of 15. Finally, we

show that the thermal stability of Al/TiN bilayers to interfa-

cial breakdown is increased from 450 to 610 C through the

use of fully dense TiN layers with complete 111 texture.

FIG. 10. XRD diffracted intensity contour maps plotted as a function of

temperature during thermal ramping of a series i and b series-ii Al/TiN

bilayers from 100 to 650 C at 3 C s1.



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ACKNOWLEDGMENTS

The authors gratefully acknowledge the financial support

of the NSF/DARPA VIP Program and the Department of

Energy under Contract No. DEAC0276ER01198 during the

course of this research. We also appreciate the use of the

Brookhaven National Synchrotron Light Source beamline

X-20C together with the expert technical assistance of Dr.

C. Lavoie and Dr. C. Cabral, Jr. of the IBM T.J. Watson

Research Center, and the facilities in the Center for Mi-

croanalysis, which is partially supported by DOE, at the Uni-

versity of Illinois.

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