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 601

With slightly higher oxygen concentration levels (J

∼ 10

−2

), coarsening during

coalescence is severely suppressed, resulting in grains with random orientation. The

competitive growth which follows is governed by anisotropic crystallographic effects [118]:

O segregates fastest at 111 surfaces. Oxygen is incorporated into the lattice of 001 and 110

crystal faces, while an oxide layer is formed on the 111 faces [119]. Oxygen tends to

accumulate at step edges on 111 surfaces, blocking step motion and leading to step bunching.

These pinning sites serve to nucleate the oxide phase. Neighboring 111 grains have rounded

edges due to oxygen segregation while 111-oriented grains in contact with 001-facets remain

sharp, as oxygen is incorporated in the latter. The 111 grains eventually develop rounded

surfaces, indicating that local epitaxial Al growth has been interrupted by an oxide layer,

above which renucleation of metal islands takes place. Crystal growth on 001-oriented grains

is unimpeded by oxygen; these grains protrude above the average ﬁlm surface and eventually

win in competitive growth. They develop the shape of truncated octahedrons bounded by a 001

top face and 111 side faces. The degree of 001 preferred orientation increases with ﬁlm

thickness and is accompanied by greater surface roughness with increasing oxygen

concentrations (Figure 12.38c).

At still higher oxygen concentrations (J

∼ 0.1–1), the oxide layer completely covers

islands of all orientations at an early stage and coarsening during coalescence is blocked.

Thus, ﬁlm growth proceeds by repeated renucleation. The ﬁlm is composed of 3D equiaxed

(globular) grains with random orientation and a zone III structure (Figure 12.38d). With

increasing oxygen concentration, the grain size decreases and can reach the nanometer range.

An important byproduct of repeated nucleation and nanograin ﬁlm formation is that surface

faceting on individual columns, and the related shadowing effects, are eliminated. Thus,

nanophase ﬁlms are inherently much smoother and, as a result, denser. The presence of oxide

phases also inhibits grain boundary migration in the bulk of the ﬁlm, preventing grain

coarsening and imparting higher thermal stability. This approach has been systematically

exploited in order to synthesize superhard nanocomposite ﬁlms based on transition metal (TM)

nitrides and carbides, e.g. nc-TMN/a-Si

[120, 121], nc-TMC/a-C [122],TiN

[123,

124],TiC

[125], and TMN/Cu [126], as well as supertough Y

-stabilized ZrO

/Au

layers [127].

As the oxygen concentration is further increased, (J

∼ 2–5), the role of the oxide and

metal phases are reversed: the oxide phase nucleates ﬁrst, while Al segregates to the surface

and forms 3D islands [117, 128]. Resulting ﬁlms are composed of metallic grains dispersed in

an oxide matrix (Figure 12.38e) [129]. Such composite ﬁlms, consisting of a low-diffusivity

matrix with higher diffusivity metallic inclusions, are the basis of a class of ceramic–metallic

coatings with diverse applications: resistors [130, 131], sensors [132], solar cell elements

[133], low-friction hard coatings (e.g. TM/a-C) [134], and smart tribological coatings that

adapt to the environment [135, 136].

602 Chapter 12

At very high oxygen ﬂuxes (J

 1), the ﬁlms consist entirely of aluminum oxide, which

for room temperature growth is amorphous. T

values exceeding 800

◦

C are required for the

synthesis of the chemically and mechanically stable κ and α crystalline phases of alumina.

There has, however, been a concerted effort to achieve hard crystalline alumina using

ion-irradiation during growth at temperatures below 500

◦

C [137].

The addition of metallic elements, rather than gas-phase species, for multicomponent

microstructural modiﬁcation has been shown to both increase and decrease the average grain

size depending on the choice of materials. For example, adding a few at.% Pt [138] or Cu

[139] to Al leads, in contrast to tissue phase formation, to the nucleation of 3D islands of the

minority phase(s) on the surface of the majority phase. Here, all adspecies have high surface

mobilities. For the case of 350

A thick Al + 4 at.%Pt layers co-evaporated onto a-C

substrates at T

= 350

◦

C, TEM imaging and diffraction analyses show that the minority

phases are Al

Pt and Al

Pt. Equiaxed grains of the minority phase decorate grain boundaries

and triple points, thus signiﬁcantly decreasing grain boundary migration and grain coarsening

[140].

The addition of Sn to Al has quite the opposite effect; rather than decreasing the grain size as

with O and Pt, it acts to enhance grain growth. In a combinatorial materials science

experiment, Barna and Adamik [92] deposited a 1000

A thick Al(Sn) ﬁlm on a-C at room

temperature by co-evaporation using a geometry and relative deposition rates such that the Sn

concentration in the ﬁlm varied from near zero at the left edge of the ﬁeld of view of a

plan-view TEM image to 10 at.% at the right edge. At low Sn concentrations, the ﬁlm is

continuous with an average grain size of 30–50

A, while at higher concentrations the ﬁlm is

still in the island growth stage with the grain size enhanced by a factor greater than 5×.Sn

appears to be acting as a surfactant in increasing Al adatom surface mobility [141].

12.6.3 Ion Irradiation Effects

Films deposited at low temperatures with little or no ion irradiation tend, as noted in Section

12.6.1, to be columnar and underdense. Figure 12.39 is a typical XTEM image of such a

structure; in this case, a TiN ﬁlm grown by reactive magnetron sputtering in pure N

, with

= 20 mtorr, at T

= 350

◦

C(T

= 0.17) and R = 0.75 ␮m/h (2.1

A/s) on amorphous SiO

with no applied substrate bias [142]. The ﬁlm lattice constant was found by XRD to be equal

to that of unstrained bulk TiN. Figure 12.39 shows that the ﬁlm grain size, initially small,

increases continuously with ﬁlm thickness while the column boundaries become increasingly

more open. The self-organized zone T (see Figure 12.33) columnar microstructure forms

through random nucleation, limited coarsening during coalescence, and competitive column

growth resulting in a microstructure consistent with that predicted by the kinetic Monte Carlo

simulations and schematic diagram in Figure 12.34. The column tops are faceted owing to

kinetic roughening, which, in combination with atomic shadowing, results in deep cusps

Thin Film Nucleation, Growth, and Microstructural Evolution 603

Figure 12.39: A bright-ﬁeld XTEM micrograph of a TiN ﬁlm deposited by magnetron sputtering in

a pure N

atmosphere (20 mtorr) on amorphous SiO

at 350



C. The ion-to-Ti ﬂux ratio J

incident at the growing ﬁlm was < 1 with an ion energy E

= 20 eV. (From [142].)

between columns and open column boundaries. The individual columns, however, are dense,

indicating sufﬁcient adatom surface mobility to sustain local crystal growth.

Many of the early applications of ion irradiation during ﬁlm growth were designed to address

issues arising from low-temperature deposition such as those apparent in Figure 12.39. The

motivation was generally to increase ﬁlm density and modify grain morphology (e.g. columnar

to equiaxed) as well as to manipulate ﬁlm texture. The experiments were carried out using

plasma-based sputtering techniques by applying a negative bias to the substrate and later by

using ion beam assisted deposition (IBAD) in which a secondary ion source is aimed at the

substrate and growing ﬁlm deposited by primary ion beam sputtering or evaporation. However,

essentially all experiments employed what is today considered to be high-energy ion

bombardment (E

> 100 eV); that is, well above the bulk lattice atom displacement threshold,

giving rise to signiﬁcant residual ion-induced lattice damage. It was not until the development,

in the late 1980s and early 1990s, of gridless Hall-current ion sources [143] for IBAD

applications and tunable magnetically unbalanced magnetron sources [144] for plasma

deposition, that the ability to independently vary ion currents (or, more importantly,

ion/deposited-atom ratios J

) at the growing ﬁlm over large values while using low-energy

ions (typically ≤ 20 eV), that it became possible to control density, grain morphology,

preferred orientation, surface smoothness, and ﬁlm stress without the collateral introduction of

detectable residual defects.

12.6.3.1 High Energy, Low Ion Flux

Mattox and Kominiak [145] were among the ﬁrst to quantitatively demonstrate that the density

of ﬁlms grown under conditions leading to porous microstructures (e.g. low T

and/or high

pressures such as are typical of DC sputtering) can be increased by ion bombardment during

604 Chapter 12

Figure 12.40: (a) Density, as a function of the negative substrate bias V

,of6-␮m-thick Ta ﬁlms

deposited by dc sputtering in Ar (adapted from [145]). (b) Average grain size d and dislocation

density n

in Ag ﬁlms deposited at room temperature as a function of the average normalized ion

energy <E

> per deposited Ag atom. (Plotted from tabulated data in [146].)

deposition. The density of 6-␮m-thick polycrystalline Ta ﬁlms deposited at 300

◦

= 0.175) by dc sputtering in Ar increased continuously with the applied negative

substrate bias V

from ∼ 14.5 g cm

−3

with V

= 0 to 16.3 g cm

−3

(bulk density = 16.6 g cm

−3

)

with V

= −500 V, as shown in Figure 12.40(a).

Huang et al. [146] investigated the effect of Ar

bombardment during the growth of Ag ﬁlms

at room temperature using a UHV dual ion-beam apparatus. They observed that the void

density decreased with increasing ion energy, consistent with the Ta dc sputtering result.

However, they also found that the average grain size d decreased, from 420 to 145

A, while the

dislocation number density n

increased from 7 × 10

to 1.3 × 10

−2

, as a function of the

normalized ion irradiation energy <E

>, deﬁned as the total Ar

ion energy delivered to the ﬁlm

per deposited Ag atom. <E

> was varied from thermal (obtained by evaporation with no ion

bombardment) to 190 eV/atom (see Figure 12.40b). In addition, the degree of 111 preferred

orientation, determined by XRD, decreased and the plane stress reversed from 0.06 GPa tensile

for underdense evaporated ﬁlms to −0.45 GPa compressive with <E

> > 42 eV/atom.

Densiﬁcation via high-energy ion irradiation at the cost of residual defect production is clearly

illustrated by the plan-view TEM micrographs in Figure 12.41 of polycrystalline Ti

0.5

ﬁlms deposited by reactive magnetron sputtering in mixed Ar/N

discharges with

tot

= 8.2 mtorr (P

= 2.9 mtorr) on stainless-steel substrates at ∼ 450

◦

C(T

∼ 0.20) with

∼ 0.9 and R = 0.2 ␮m/minute (33.3

A/s) [147]. The applied negative substrate bias V

ranged from 0 to 250 V (corresponding to ion energies of ∼ 6.5 to 255 eV when the plasma

potential is included). XTEM images (not shown) reveal that ﬁlms grown with V

= 0 have a

columnar microstructure with intercolumnar void networks similar to those in Figure 12.39.

Film porosity decreases sharply with V

> 100 eV until, with V

> 120 V, no voids are observed

Thin Film Nucleation, Growth, and Microstructural Evolution 605

Figure 12.41: Transmission electron micrographs of polycrystalline Ti

0.5

N ﬁlms deposited by

reactive magnetron sputter deposition on stainless-steel substrates at T

∼ 450



C as a function

of the applied negative substrate bias V

with J

∼ 0.9. (a) V

= 0; (b) 75 eV, (c) 120 eV, and

(d) 250 eV. (From [147].)

using underfocus and overfocus contrast. However, the ﬁlm compressive stress (owing to a

combination of ion trapping and recoil implantation) increases and high dislocation loop

densities are observed within individual grains in panels (b)–(d). In fact, the defect densities

are large enough to disrupt column growth and force renucleation, as evidenced by the

presence of Moir

e fringes, leading to smaller grain sizes.

X-ray and electron diffraction patterns also reveal a very slow change in TiN preferred

orientation from 111 to 002 with increasing E

. However, the ion energy required to complete

the transition is > 800 eV, for which the ﬁlms have unacceptably high stress levels [148]. Thus,

the use of high-energy, low-ﬂux ion irradiation is not a practical approach for controlling ﬁlm

texture. The formation of 002 texture under such conditions is directly related to collision

cascade effects [149]. Grains with open channel directions, such as 002, have higher survival

probabilities due to the anisotropy in collision cascades; that is, the ion energy is distributed

over larger depths in open channels leading to lower sputtering yields and less lattice distortion.

In summary, densiﬁcation obtained in the high-energy ion irradiation regime (E

> 100 eV)

comes at a steep price, leading to high defect densities [150], high compressive stresses

[151–153], and inert gas incorporation [148]. Ar concentrations C

in TiN layers deposited on

amorphous SiO

at 350

◦

C as a function of E

, with J

≤ 1, are below 0.5 at.% with

606 Chapter 12

< 100 eV, while at higher ion energies, C

increases approximately linearly from 1 at.% at

200 eV to > 3 at.% at 800 eV.

12.6.3.2 Low Energy, High Ion Flux

Magnetron sputtering is a primary technique for the growth of polycrystalline thin ﬁlms. It

provides high deposition rates due to the Lorentz force, arising from the cross-product of the

applied electric ﬁeld used to accelerate the ions to the target and the permanent magnetic ﬁeld,

which traps the plasma immediately adjacent to the target and efﬁciently utilizes secondary

electrons in the discharge to produce new ions. However, this very advantage becomes a

disadvantage when attempting to use ion bombardment at the substrate for microstructural and

compositional modiﬁcation during ﬁlm growth. The strong electromagnetic trap near the target

means that the substrate and growing ﬁlm are essentially outside the plasma; thus, a high

substrate potential is required to obtain signiﬁcant ion bombardment and, even then, the ion

ﬂux at the substrate is small. For typical magnetron discharges, depending on the target

material, gas, pressure, and system geometry, ion/deposited-atom ratios at the substrate with

= −100 V are of the order J

∼ 0.1–0.5. This dramatically changed with the

development of the tunable unbalanced magnetron which incorporates a Helmholtz coil

external to the permanent magnets of the magnetron to unbalance the magnetic circuit and

controllably open a leak in the plasma [144]. High ion/deposited-atom ratios, J

> 50, can

be achieved using very low ion energies (E

< 10–20 eV) with no signiﬁcant change in ﬁlm

deposition rates. Moreover, the ion energy and ion ﬂux incident at the substrate are

independently controlled by V

and the external magnetic ﬁeld B

ext

, respectively.

≤ 20 eV is below the threshold for bulk lattice atom displacement in NaCl-structure TM

nitrides. The residual stresses remain low in as-deposited ﬁlms, yet the effects on texture and

microstructure are dramatic. In some of the earliest experiments using unbalanced magnetrons

[144, 154, 155], it was clearly demonstrated that the average ion bombardment energy per

deposited atom <E

> is not a universal parameter for microstructural and texture modiﬁcation

as had been proposed by many authors. Figure 12.41 conclusively shows, for example, that

changes in the stress state, stoichiometry, and texture of NaCl-structure polycrystalline

0.5

N ﬁlms deposited on a-SiO

follow completely different reaction paths depending on

whether <E

> is varied by changing E

(with a constant value of J

) or varying J

(with

constant E

) [154, 155]. Similar results have been reported for TiN [142],TaN[113], ScN

[156], and many other materials.

The polycrystalline 1-␮m-thick NaCl-structure Ti

0.5

N ﬁlms used to obtain the

experimental results summarized in Figure 12.42 were deposited by reactive magnetron

sputtering from a TiAl alloy target on a-SiO

at 250

◦

C in pure N

discharges at 20 mtorr

[155]. The relatively high pressures were chosen in order to thermalize the sputtered metal

atoms and neutralized ions reﬂected from the target such that the only energetic species

incident at the growing ﬁlm were accelerated nitrogen ions (∼ 96.5% N

and 3.5% N

) [67]

Thin Film Nucleation, Growth, and Microstructural Evolution 607

Figure 12.42: (a) Normalized X-ray diffraction peak intensities I; (b) composition ratios

N/(Ti + Al); and (c) out-of-plane lattice parameters a

from Ti

0.5

N alloys deposited on

amorphous SiO

at 250



C under N

ion irradiation as a function of the normalized energy <E

per deposited metal atom varied through changes in either the ion energy E

or the ion/metal ﬂux

ratio J

. (Adapted from [155].)

whose ﬂux was controlled by varying B

ext

. The results in Figure 12.42 show that raising <E

through increases in E

, with J

maintained constant at 1, increases the out-of-plane lattice

constant a

(panel c) and leads to trapping of excess N in the ﬁlm (panel b). Both of these

effects are signatures of a rapidly increasing compressive stress. TEM and XTEM analyses

show that ﬁlms grown with E

=20eV(∼ 10 eV per N following collisionally induced

dissociation) and J

= 1 are underdense with both intercolumnar and intracolumnar

porosity and an average column size d = 300 ± 150

A. Increasing E

leads to densiﬁcation and

a reduction in column size to ∼ 250

A at 85 eV with a small increase in the volume fraction of

002 grains to ∼ 15%. However, at E

= 100 eV, the ﬁlms are no longer single phase and instead

608 Chapter 12

consist of a TiN-rich columnar structure, with d ∼ 150

A, and wurtzite-structure AlN-rich

precipitates of diameter ∼ 30

A. Thus, increasing the ion energy in order to achieve

densiﬁcation in low-temperature Ti

0.5

N ﬁlms results in residual lattice defects which

create compressive stress and, at E

> 85 eV, phase separation.

In sharp contrast, maintaining E

constant at 20 eV and increasing <E

> through increases in

still leads to densiﬁcation, but with no signiﬁcant ion-induced stress, and the layers

remain stoichiometric and singe phase. The column size distribution becomes more uniform

with no measurable change in column size. An additional major beneﬁt of this approach to

coupling kinetic energy to the growing ﬁlm during low-temperature deposition is shown in the

upper panel of Figure 12.42. That is, the use of low-energy, variable ion ﬂux irradiation allows

ﬁlm preferred orientation to be set at any value between 111 and essentially complete 002.

This is important in many applications due to the anisotropic properties of TM nitrides. The

mechanism for the change in texture, which involves a combination of ion-irradiation-induced

surface chemistry, anisotropic adatom diffusion dynamics, and momentum transfer, is

discussed in the following sub-section.

Ion-irradiation-induced surface smoothening

There are many polycrystalline thin ﬁlm applications in which surface and interface

smoothness is crucial to device operation. Examples include metal multilayer X-ray mirrors

for semiconductor device lithography and giant magnetoresistance (GMR) devices such as

non-volatile magnetic access memory and read head sensors in hard disk drives. The GMR

effect derives from changes in electron-spin-dependent resistance across multilayer stacks,

consisting of high-conductivity ﬁlms sandwiched between ferromagnetic layers, owing to

orientation switching of an external magnetic ﬁeld. If the ferromagnetic layers have a

roughness comparable in amplitude and wavelength to the spacer layer thickness, typically

∼ 10–20

A, it is difﬁcult to reverse the magnetic ﬁeld of the ferromagnetic layers because of

eel coupling. In addition, layer intermixing increases spin-independent scattering and loss of

local magnetic alignment.

GMR multilayers are generally composed of material pairs with large miscibility gaps, and

deposition is carried out at low temperatures, in order to minimize intermixing by thermal

diffusion during deposition. The best GMR layers are synthesized using low-energy ion

irradiation for reasons illustrated in Figure 12.43 showing molecular dynamics simulations,

using embedded atom potentials, of two three-layer Cu/Co/Cu stacks. The experiments were

designed to probe kinetic effects during the asymmetric growth of Cu/Co versus Co/Cu

interfaces [157]. Both trilayer stacks were grown with thermal atoms (evaporation), but

=6eVAr

ion irradiation with J

= 5 was added in panel (b). (Note that E

is of the same

order as the average energy of ballistic sputtered atoms, 5–10 eV, incident at the ﬁlm growth

surface during magnetron sputter deposition at low pressures, ∼ 1 mtorr, with a

target-to-substrate separation of a few cm.) The bottom Cu layer consists of eight planes of

Thin Film Nucleation, Growth, and Microstructural Evolution 609

Figure 12.43: Molecular dynamics simulations of the growth of Cu/Co/Cu trilayers by sequential

deposition of Co and Cu overlayers at room temperature on a Cu(111) substrate consisting of

eight layers. The metal beam ﬂuxes are thermal in both (a) and (b), but E

=6eV Ar

ion

irradiation, with J

= 5, is added in (b). (Adapted from [157].)

Cu(111); the Co and Cu overlayers were deposited at room temperature. In the thermal

deposition case, panel (a), the Co diffusivity is low leading to a rough Co surface and Co/Cu

interface. Thus, as the upper Cu layer buries the irregular Co surface, an even rougher upper

surface is formed.

Panel (b) corresponds to the growth of the same trilayer system, but this time with very low

energy ion irradiation simulating ion-assisted evaporation. The experiment is also a crude

simulation of ballistic sputter deposition since ∼ 83% of the arriving particles are energetic

and the mass mismatch among Ar, Cu, and Co is not large. Ion-irradiation-enhanced adatom

610 Chapter 12

Figure 12.44: Epitaxial thickness h

epi

of Si/Si(001) layers grown with hyperthermal E

=11eV

beams () and with MBE (䊉)atR ∼ 1

A/s as a function of T

. (Hyperthermal beams from [160];

MBE from [161].)

mobilities greatly decreases the roughness of both interfaces, and the surface, without adding

to interfacial mixing. Clearly, the driving effect here is momentum transfer from the fast

incident particles to surface atoms. Thus, the ion mass, ion energy, ion-to-thermal ﬂux ratio,

and ion angle of incidence can all be tuned to affect adatom mobilities and interfacial mixing

in slightly different ways [10, 158, 159].

Experimentally, Lee et al. [160] have shown (Figure 12.44) that the epitaxial thickness h

epi

low-temperature (T

= 80–300

◦

C, T

= 0.21–0.34) homoepitaxial Si/Si(001) layers grown

from hyperthermal Si beams is increased by up to an order of magnitude above that recorded

for MBE Si(001) epitaxy [161] at approximately the same deposition rate, R ∼ 1

A/s (Figure

12.44). The hyperthermal Si atoms were incident at the growing ﬁlm with an orthogonal

energy component of E

= 11 eV. In both cases, continued growth at ﬁlm thicknesses h > h

epi

leads to increased kinetic surface roughening, growth front breakdown, and loss of epitaxy

with a non-reversible transition to amorphous layer deposition. h

epi

ranges from ∼ 100

Aat

= 100

◦

C, to 500

Aat150

◦

C, to > 1 ␮mat300

◦

C for Si(001) layers grown from

hyperthermal species. For comparison, h

epi

for MBE Si(001) varies from 70

AatT

= 150

◦

Cto

1200

AatT

= 300

◦

C for MBE Si(001).

The slope change in the h

epi

versus T

curve for hyperthermal species in Figure 12.44 is

explained as follows. While the energy of the hyperthermal growth species is too low to give

rise to detectable residual bulk defect densities, they collisionally dissociate surface dimers

and small clusters, thus resulting in higher adatom densities on all exposed surface levels. At