Martin P.M. Handbook of Deposition Technologies for Films and Coatings, Third Edition: Science, Applications and Technology

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Thin Film Nucleation, Growth, and Microstructural Evolution 611

sufﬁciently low growth temperatures (< 225

◦

C for this set of deposition conditions), this leads,

on a given level, to an increased 2D gas supersaturation which in turn results in higher island

nucleation rates with smaller average cluster sizes and, therefore, enhanced interlayer mass

transport rates and epitaxial thicknesses. This effect is less important at higher growth

temperatures as thermal diffusivities overcome the excess supersaturation due to collisionally

formed adatoms.

Further increases in h

epi

in the very low temperature regime, < 225

◦

C, can be obtained by

manipulating island nucleation kinetics in another manner, through changes in R. Initial

decreases in R from 1.0 to 0.6

A/s result in h

epi

increasing from 500 to 800

AatT

= 150

◦

owing to increased adatom diffusion lengths at lower R values giving rise to higher step

crossing probabilities and decreased surface roughness. However, there is a competing

reaction. Increasing adatom diffusion distances by decreasing R at constant T

also leads to

increased island sizes and a reduction in the effects of energetic condensation on nucleation

kinetics, both of which tend to increase surface roughness. Thus, for a given T

, there is a

maximum in h

epi

vs R and the epitaxial thickness decreases again at even lower deposition

rates.

Ion-irradiation-induced preferred orientation

As shown in Figure 12.42, the texture of polycrystalline Ti

0.5

N ﬁlms (as well as other

NaCl-structure TM nitrides) can be controlled from nearly complete 111 to 002 by varying

, while maintaining E

≤ 20 eV, during low-temperature ﬁlm growth. 002 is the

low-energy surface for NaCl-structure TM nitrides [114, 162] and the ﬁlms are not under

signiﬁcant strain. Thus, the 111 preferred orientation observed with low J

ratios is

kinetically limited. While 111 and 002 grains, approximately 50% each by area, coexist during

the early stages of ﬁlm growth, the 111 grains gradually overgrow the 002 oriented grains until

at thicknesses > 1500

A, the ﬁlm has a nearly complete 111 texture [155] similar to the kinetic

Monte Carlo simulation in Figure 12.34 (with, in the present case, the darker contrast regions

representing 002 grains and the lighter regions 111 grains) and the XTEM micrograph in

Figure 12.39. (See [113, 142] for similar results during TiN and TaN growth as a function of

.) This is due to a combination of ion-induced surface chemistry, anisotropies in surface

diffusivities, and differences in adatom potential energies.

For the growth of TM nitrides by reactive sputter deposition, the interplay and competition

among several interrelated surface reaction and diffusional processes determine ﬁlm growth

kinetics, surface morphology, ﬁlm microstructure, and texture [114]. The rate-limiting step for

TM nitride ﬁlm growth is cation incorporation which, in these experiments, proceeds at a rate

of several ML/s. The N

supply rate is much larger, > 10

ML/s. If N

molecules reach Ti

adatoms not yet incorporated in the lattice, they undergo dissociative chemisorption. However,

all other incident N

molecules are only physisorbed and desorb at kinetic rates with estimated

lifetimes of ≤ 10

−9

s. This results in (002) surfaces which exhibit an essentially

612 Chapter 12

bulk-terminated structure with equal cation and anion number densities while polar (111)

surfaces are fully N-terminated [114]. Ti adatoms form one N backbond on (002) and three N

backbonds on (111) surfaces. Thus, cation diffusivities and potential energies are higher on

001-oriented grains than on 111 grains and, following the arguments in Section 12.6.1 for

competitive grain growth, Ti adatoms have a larger chance of becoming trapped at 111 sites.

Consequently, 111 grains slowly and inexorably expand at the expense of the 002 grains under

low-T

, low-ion-ﬂux growth conditions.

ω-2θ and glancing-angle XRD scans from 1 ␮m thick low-temperature TiN (T

= 350

◦

∼ 0.17) and Ti

0.5

N(T

= 250

◦

C, T

∼ 0.15) layers grown in pure N

20 mtorr on a-SiO

with E

=20eVandJ

values ranging from ∼ 1 to 15 contain (111)

and/or (002) reﬂections (no other orientations were detected). (See [113] for similar results

for TaN.) Normalized 111 and 002 peak intensities are plotted as a function of J

Figure 12.45. Films grown with high ﬂux ratios (J

≥ 5) exhibit an 002 texture, while those

grown with low J

have a 111 preferred orientation. XRD pole ﬁgures reveal, in addition,

that all diffraction peak maxima are aligned along the growth direction and that peak

intensities are azimuthally symmetric. Measured residual stresses are low at all J

values.

XRD analyses as a function of ﬁlm thickness h and J

, as well as cross-sectional

selected-area electron diffraction results from thick ﬁlms, reveal competitive grain growth up

to h ∼ 1500

A, with both 111 and 002 grains present for J

≤ 10. At J

= 1, the ﬁlms are

porous and evolve to pure 111, while at J

∼ 5–10 the ﬁlms are fully dense and it is the 002

columns which evolve by overgrowing the 111 columns. In addition, the 002 layers exhibit

much smoother surfaces than the faceted 111 layers.

Figure 12.45: Normalized 111 and 002 peak intensities I from 1-␮m-thick TiN [142] and

0.5

N [155] layers deposited by reactive magnetron sputtering in pure N

(20 mtorr) on

amorphous SiO

substrates at T

= 350



C (TiN) and 250



C(Ti

0.5

N) with E

=20eVasa

function of J

Thin Film Nucleation, Growth, and Microstructural Evolution 613

The above behavior can be understood based on the arguments, presented in Section 12.6.1,

regarding competitive growth between grains exposing low- versus high-diffusivity planes to

the growth front with low-energy high-ﬂux N

ion irradiation switching the balance between

111 and 002 orientations. Incident collisionally dissociated 20 eV N

ions provide a

continuous source of atomic N which can chemisorb on 001 grains ([001] is a non-polar

direction in the NaCl-structure), but not on N-terminated 111 grains. This is in contrast to

incident thermal N

species which, as discussed above, have very short lifetimes on both

surfaces in the absence of a local population of free cations. Thus, increasing J

corresponds to raising the steady-state N coverage θ

on 001 grains while θ

remains at a

constant N-terminated value on 111 grains. This has the important consequence of decreasing

cation mean free paths on (002) surfaces due to capture of N adatoms to form TiN

(i = 1–3)

admolecules, in the case of TiN, which have much lower surface diffusivities than Ti adatoms

as quantitatively demonstrated by TiN(001) nucleation results in Section 12.4.1 [62, 68] and

ab initio density functional calculations [114]. That is, in the presence of an atomic N source,

the 002 chemical potential decreases below that of 111. Consequently, the net ﬂux of cations

from 002 to 111 oriented grains is reversed under low-energy, high-ﬂux N

irradiation

conditions, resulting in the development of 002 texture.

In addition to controllably switching the texture from 111 to 002 with increasing ion ﬂux, there

is a corresponding increase in layer density and a decrease in surface roughness from

underdense layers, with intercolumnar voids and self-organized growth mounds separated by

deep surface trenches, to fully dense layers with smoother surfaces. The densiﬁcation arises

from less pronounced kinetic roughing due to ion-irradiation enhanced adatom surface

mobilities resulting in smoother surfaces with less atomic shadowing.

The above results demonstrated that increasing J

with E

= 20 eV during growth of

NaCl-structure TM nitride ﬁlms results in a transition from underdense 111 to a dense 002

preferred orientation. With the introduction of texture inheritance, however, it becomes

possible to select the preferred orientation during the nucleation stage. This concept has been

used to achieve dense, highly 111-oriented TiN ﬁlms [163]. (Similar results were obtained

with TaN [164].) Ti underlayers, 250

A thick, are ﬁrst deposited on amorphous substrates at

◦

C(T

= 0.18) with J

= 2 and E

= 11 eV. The Ti layers are fully dense with a zone II

structure and very strong 0002 orientation. The TiN overlayers are then deposited with high

values (> 10) and E

∼ 20 eV, which on bare amorphous substrates results in dense 002

layers. In this case, however, local epitaxy on 0002 Ti grains yields fully dense TiN layers with

a strong 111 texture; competitive grain growth is completely eliminated. Texture inheritance

selects the preferred orientation while low-energy ion irradiation provides densiﬁcation.

In addition to out-of-plane texture, in-plane preferred orientation can also be controlled using

ion irradiation by taking advantage of anisotropic collisional cascade effects [165]. However,

experiments to date have primarily been carried out with high-energy ions in an IBAD

614 Chapter 12

conﬁguration. Selectively oriented layers of refractory metals [166], oxides [167, 168], and

nitrides [169] have been obtained and used as templates to provide pseudomorphic texture

inheritance to low-temperature overlayers. An example is the growth of highly 002 textured

polycrystalline MgO ﬁlms for seed layers on which to deposit oriented oxide superconductors.

The strategy is to align an ion beam such that it is orthogonal to a high-symmetry, low-density

plane in a growing MgO crystallite oriented with the (002) plane parallel to the substrate

surface. Ion energy in open channel directions is distributed over larger depths yielding less

lattice distortion and lower sputter yields. Therefore planes perpendicular to an open channel

direction have higher survival probability. Since MgO has the NaCl structure, the beam is

aligned orthogonal to the (101) plane at 45

◦

to the [001] direction corresponding to the

selection of in-plane preferred orientation along [100] [168].

Three-dimensional nanostructures

Three-dimensional nanostructures have recently become of interest owing to the possibility of

achieving superhard layers with high wear resistance [121, 170], layers which actually

increase in hardness with increasing temperature [124] (important for high-speed machining,

turbine blade coatings, etc.), and supertough layers [171]. A general strategy is evolving which

combines segregation of incommensurate phases to force periodic renucleation, during

low-temperature deposition, with low-energy ion irradiation to control both phase distribution

and preferred orientation [123, 124, 170, 171].

Relatively simple examples are the self-organized Ti

0.8

0.2

N nanostructures shown in

Figure 12.46 [171]. They are grown on amorphous SiO

(R = 0.63

A/s) by reactive magnetron

sputtering in pure N

at 5 mtorr with T

= 350

◦

C. The constituent TM and rare-earth nitrides

TiN (a

TiN

= 4.240

A) and CeN (a

CeN

= 5.021

A) exhibit a deep miscibility gap with a large

lattice parameter mismatch. Z-contrast scanning XTEM, nanodiffraction, and electron energy

loss spectroscopy analyses, combined with the high-resolution TEM and XTEM lattice images

in Figure 12.46, show that the nanostructures are fully dense with a nearly complete 002

texture. The upper nanocomposite consists of ∼ 20

A diameter nanospheres encapsulated in a

fully percolated disordered CeN-rich tissue phase. Continuous CeN surface segregation during

growth of TiN-rich nanocrystallites leads to renucleation. TiN then segregates out of the

growing CeN phase until it reaches a critical supersaturation to nucleate a new nanocrystallite.

Increasing the ion irradiation intensity during TiCeN deposition results in the formation of

quasi-periodic, 20

A wide, CeN-encapsulated TiN-rich nanocolumns (lower image in

Figure 12.46). Ion-irradiation enhanced surface mobilities combined with a large strain energy

driving force gives rise to a surface-initiated spinodal decomposition process allowing the

continued formation of adjacent CeN-rich and TiN-rich nanocolumns. The in-plane periodicity

can be tuned by changing the overall Ce/Ti ratio, the ion irradiation conditions, T

, and (to a

lesser extent) R.

Thin Film Nucleation, Growth, and Microstructural Evolution 615

Figure 12.46: High-resolution XTEM micrographs, with corresponding selected-area diffraction

patterns, from 3D Ti

1−x

N nanostructures, of average composition x = 0.20, grown by reactive

magnetron sputtering in pure N

(5 mtorr) on amorphous SiO

at T

= 350



C under different ion

irradiation conditions. The nanocomposites consist of approximately spherical TiN-rich

nanocrystallites (upper image) and nanocolumns (lower image), both ∼ 20

A in size, encapsulated

in a disordered CeN-rich phase. (Adapted from [171].) The dotted lines and labels are to guide

the eye.

A related example with a more complex nanostructure has the overall composition TiB

2.4

[123]. Since hexagonal-structure TiB

is a line compound, the excess B does not remain in

solution. The nanocomposite is self-organized at two length scales and consists of

quaisperiodic 200-

A-wide stoichiometric TiB

columns separated by 1–2 ML of B. The 200

nanocolumns are, in turn, composed of even smaller 50-

A-wide nanocolumns, also separated

by a B tissue phase. The structures have a very high hardness since the narrow columns inhibit

the nucleation and glide of dislocations while the strong covalent bonding of the B tissue phase

prevents grain boundary sliding. Going a step further to synthesize Ti–B–N with overall

compositions, such as TiB

0.80

0.83

(38 at.% Ti, 30.5 at.% B, 31.5 at.% N), within a deep

miscibility gap in the TiN–TiB

–BN ternary phase ﬁeld, results in self-organized

nanostructures which are not only hard, but whose hardness increases with temperature owing

to increased ordering in the covalent tissue phase [124].

616 Chapter 12

12.7 Conclusions

Atomic scale understanding of microstructural evolution, necessary for the controlled

manipulation of the properties of thin ﬁlm systems, is increasing rapidly. However, we are still

far from having accurate quantitative descriptions. The ﬁeld is presently driven both by

scientiﬁc curiosity arising from the richness of surface and bulk phenomena involved and,

more and more, by technological demands to provide materials with properties that do not

exist at present. Rapid progress is now possible owing to the availability of at least ﬁrst-order

atomic-scale design rules. Overall, the ﬁeld of thin ﬁlms continues to be vital, intellectually

challenging, and a long way from mature as it moves toward the ultimate goal of being able to

develop designer materials, at the atomic scale, possessing any desired set of properties.

Acknowledgements

The author gratefully acknowledges the ﬁnancial support of the Department of Energy,

Division of Materials Science. He has also beneﬁted from many years of close collaboration

and friendship with Ivan Petrov (University of Illinois) as well as continuous discussions and

scientiﬁc interactions with Lars Hultman (Link

oping University) and Peter Barna (Hungarian

Academy of Science).

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