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This is to certify that the

thesis entitled

EFFECT OF PLASMA ENERGY, PLASMA POWER, WORKING DISTANCE, AND
INERT GAS PRESSURE ON COMPOSITION OF SPUTTERED THERMOELASTIC

[)ate

TITANIUM-NICKEL THIN FILMS
presented by

Jiun-Chung Lee

has been accepted towards fulfillment
of the requirements for

Master ' 5 Materials SEIénce

degree in

 

 

 

 

Major professor

August 17, 1994

 

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DATE DUE DATE DUE DATE DUE

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

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«term

”3-9.1

EFFECT OF PLASMA ENERGY, PLASMA POWER, WORKING DISTANCE, AND
INERT GAS PRES SURE ON COMPOSITION OF SPUTTERED THERMOELASTIC
TITANIUM-NICKEL THIN FILMS

By
Jinn-Chung Lee

A THESIS
Submitted to
Michigan State University
in partial fulﬁllment of the requirements
for the degree of

MASTER OF SCIENCE

Department of Materials Science and Mechanics

1994

ABSTRACT

EFFECT OF PLASMA ENERGY, PLASMA POWER, WORKING DISTANCE, AND
INERT GAS PRES SURE ON COMPOSITION OF SPUTTERED THERMOELASTIC
TITANIUM-NICKEL THIN FILMS

By

Jinn-Chung Lee

Thin ﬁlms of TiN i are of interest for application to shape-memory microactuators.
The transformation temperature of TiN i shape memory alloys is extremely sensitive to the
alloy composition. Increasing the Ni content beyond stoichiometry decreases the
transformation temperature by approximately 100 K per 1 at% Ni. Therefore, to maintain
reasonable control over transformation temperature, it is desirable to control ﬁlm
composition to within 1 0.1 at%. The present study has focused on the effect of sputter
deposition process variables on composition in TiN i thin ﬁlms.

By adjusting the process variables, thin ﬁlms with compositions showing positive
and negative deviations from the target-alloy composition have been obtained. It was
found that the plasma power had a critical effect on the ﬁlms composition, although the
plasma voltage and gas pressure also affected it. The results are understood in terms of the
sputtering yield of alloy constituent elements and the mean free path of energetic neutral

atoms in the plasma.

ACKNOWLEDGMENTS

I would like to give thanks to several important people.

To Dr. Grumman, my major professor, who through his knowledge and guidance made

this study successful.

To Liuwen Chang for his initiative on the study of sputter-deposited TiN i ﬁlm

composition.

To Seungho Nam for his constant help on the experimental technology.

To Zhiwei Cai for the patient instruction on the use of the SEM and composition analysis.
To Li Ho for his effort on the maintenance'of the sputtering system.

To Alan Gibon for his help on the thesis writing.

And to my parents, their constant support and substantial ﬁnancial contributions, made my

education possible.

TABLE OF CONTENTS

LIST OF TABLES
LIST OF FIGURES

CHAPTER 1 INTRODUCTION

CHAPTER 2 LITERATURE REVIEW

2.1 TiN i Shape Memory Alloys
2.1.1 Shape Memory Properties
2.1.2 Phase Stability and Occurrence in TiNi Alloys
2.1.3 The Effects of Alloy Composition on
The Transformation Temperature

2.2 Mechanical and Metallurgical Studies of TiNi Thin Films
2.3 Deposition by Sputtering

2.4 Principles of Sputtering
2.4.1 Properties of Glow Discharge
2.4.2 Physics of Sputtering
2.4.3 Sputter Deposition Film Growth
2.4.4 Magnetron Sputtering
2.4.5 Deposition of Alloys Films

2.5 Composition Analysis
2.5.1 Principle of EDS Analysis
2.5.2 Characteristic X-ray Intensity
2.5.3 ZAF Correction

CHAPTER 3 EXPERIMENTAL METHODS
3.1 Materials
3.2 Deposition Equipments
3.2.1 Sputtering Geometry
3.2.2 Pumping Systems and Gas Pressure Monitors

3.2.3 Magnetically Sputtering Cathode Assembly and Power Supply
3.2.4 Thickness Monitor

iv

3.3 Deposition Procedure 54

3.3.1 Preparation 54

3. 3. 2 Processing 55

3.4 Composition Analysis 56

3.5 Thickness Measurement 57

CHAPTER 4 RESULTS AND DISCUSSION 58
4.1 Factors Determining The Composition of

The Sputter-Deposited Thin Films 58

4.1.1 Preferential Sputtering and Component Sputtering Yield Ratio 59

4. 1.2 Angular Distribution of Sputtered Atoms 62

4.1.3 The Development of Target Surface Topography 66

4.1.4 Resputtering of Growing Thin Films 72

4.1.5 Planar Magnetron Source 76

4.2 Effects of Process Parameters on Thin Film Composition 77

4.2.1 Effect of Target Erosion 79

4.2.2 Effect of Plasma Voltage 84

4.2.3 Effect of Plasma Power 87

4.2.4 Effect of Gas Pressure 93

4.2.5 Effect of Working Distance 97

CHAPTER 5 CONCLUSIONS 99

BIBLIOGRAPHY 101

Table

2.2:
2-3:
4-1:
4—2:
4-3:
4-4:

4-5:

4-7:

4-8:

LIST OF TABLES

Threshold Energies (eV) [28]

Sputtering Conditions and Film Compositions in [31]

Sputtering Yield for Metals in Argon (atoms/ion) [36]

Properties Comparison of Titanium and Nickel

Histories of Targets in the Different Factor Tests (W - hour)

Process Condition for Target Erosion Test and the Compositional Results

Process Conditions and the Compositional Results in the Voltage Effect
Test

Process Conditions and the Compositional Results in the Power Effect
Test

Process Conditions and the Compositional Results in the Pressure Effect
Test

The Mean Free Path of Ar, T1, and Ni at Pressure of 0.44Pa, 0.64 Pa, and
1.01 Pa

Process Conditions and the Compositional Results in the Working
Distance Effect Test

Page

12
16
26
61
78

85

88

94

95

98

LIST OF FIGURES

Figure Page
2- 1: The shape memory effect is described with reference to a plot of electrical
resistance vs. temperature from which the characteristic transformation 5
temperature Ms, Mf, As and Af are determined [7].
2-2: Equilibrium Phase Diagram for TiN i Alloys [8] 6
2-3: Proposed equilibrium diagram of the 'li-Ni binary system in the vicinity
of the equiatomic composition at low temperature [9]. 6
2-4: Isothermal Cross Section Through the Ti-Ni-O Phase Diagram at 1200 K
[14] 7
2-5: The Dependence of the Transformation Temperature (M,) on Nickel
Content in TiN i Alloys [6] 9
2-6: The change of Ti concentration of the IBAD processed TiNi ﬁlms as a
function of HA ratio, relative to the expected composition of the ﬁlm in 13
the absence of an assist-beam. Data are shown for 50, 100, and 500 eV
beam [26]
2—7: The Valve Design by Using TiNi Thin Films [3] 15
2-8: The Mirror Actuator Concept by Using TiNi Thin Films in [3] 15
2-9: [34] .
(a) Structure of a Glow Discharge in a DC. Diode System 19

(b) Charged Particle Concentration in a Glow Discharge
(0) Voltage Variation in a DC. Diode Glow Discharge

2-10: Interactions of Ions with the Surface [34] 23
2-11: (a) Binary Collision Between Atom A and B, Followed by a Binary
Collision Between Atom B and C (b) Collision Process Responsible for 24
Sputtering and Fast Neutral Generation [36]

2- 12: Variation of Sputtering Yield with Angle of Incidence for 1 keV Ar Ion 25
Incident on Ag, Ta, Ti and Al [37]

2- 13: Sputtering Yield of Nickel as a Function of Ion Energy and Ion Mass [27] 27

2-14:
2-15:
2-16:

2-17:
2-18:

2-19:

3-1:
3-2:
4-1:

4-8:
4-9:

4-10:
4-11:
4-12:

4-13:

Formation of Thin Film [38]

Particles mearding the Substrate in Sputter-Deposition [34]

Motion of an Electron Ejected from a Surface with Velocity v into a
Region of Magnettic Field B Parallel to the Surface [34]:

(a) with No Electric Field

(b) with a Lineme Decreasing Field

Schematic Drawing of Planar Magnetron Target and its Cross Section [34]

Schematic Distribution of Electrons and Activated Volume in a Thick and
Thin Specimen [45]

Schematic Diagram of X-Ray Generation in Sample and the Absorption of
Photons from Each Depth 2 Enrout to the Detector, Depending on the X-

Ray Take-Off Angle 0 [43]
Schematic Diagram of DC. Magnetron Sputtering System
Exploded View of TORUS-2C Magnetron Sputtering Source

Angular Distribution of Sputtered Atoms from a Polycrystalline Target
[57]

The Geometric Relationship of Samples and Target Surface

An Eroded Target with its Localized Erosion Angle

The Grain Structures of Target Surface before Ion Bombardment (100 x)
The Topography of a Target Surface Sputtered for 15 Horus

The Cone Structure on the Target Surface

Schematic Diagram Showing Impurity-Induced Cone and Subsequent Pit
Formation [62]

Schematic Representation of an Unbalanced Magnetron [68]

A Typical Voltage-Current Characteristic for Planar Magnetron Cathode at
Various Pressures [69]

Composition Change as a Function of the Target Consumption
Composition Change as a Function of Plasma Voltage

Local Gas Density Measured Near the Magnetron Cathode During the
Sp8uttering of Cu in Ne, Ar, and Kr as a Function of Discharge Current
[7 ]

Composition Change as a Function of Plasma Power

32

36

37

41

45

52

63
65
65
67
68
69
7 l

74
75

80
85

89

88

4-14: Composition Change as a Function of the Gas Pressure

4-15: Composition Change as a Function of Working Distance

94
98

CHAPTER 1 INTRODUCTION

Shape memory properties in the TiNi alloys were ﬁrst discovered in the 19603 at
the Naval Ordnance Laboratory [1]. Since that time, extensive studies have been devoted
to the development and application of bulk TiNi alloys. The shape-memory effect (SME)
is due to the thermoelastic martensitic phase transformations at a given transition
temperature. Thus, the phase transformation temperatures are the most important
properties of TiN i shape memory alloys because they determine the working temperature
range of the materials. Since the transformation temperature for TiN i alloy is very
sensitive to the alloy composition [2], precise composition control in the deposition
process is essential.

TiN i alloys have been used in thin ﬁlm form as the material basis for
microactuators in microelectromechanical systems, and as surface microalloys for
structural metals [3, 4]. Usually, alloy thin ﬁlms are fabricated by a physical vapor
deposition process such as sputtering or evaporation. In the present study, TiN i thin ﬁlms
were deposited by a single-target dc magnetron sputtering system in an argon atmosphere.
Sputtering was chosen a means for producing ﬁlms because it is recognized to be the most
reliable method for control of composition [5]. However, the composition of a sputter-
deposited TiN i ﬁlm can be quite different from that of the sputtering target because of the
different process parameters, such as distance between target and substrate, gas pressure,
plasma potential (voltage), and substrate temperature. Nevertheless, the composition
change should be reproducible for a given set of conditions. Therefore, the composition
in the deposited ﬁlms are functions of the process variables.

Studies on TiNi thin ﬁlms have been reported previously. Most of them have been
concerned with the mechanical and metallurgical properties of thin ﬁlms. The objective of
the present work is to determine the working conditions under which the compositions of

TiNi thin ﬁlms can be made to show both positive and negative deviations from the TiN i

2

target alloy composition. This is desirable so that small, well-controlled deviations from
stoichiometry can be made using a single stoichiometric target which is stable with respect
to precipitation of intermetallic precipitates. The signiﬁcance of the effect of various
process variables on the composition change are surveyed. Through this study, a precise
knowledge of the relationship between thin ﬁlm composition and sputter deposition
process parameters can be obtained. Furthermore, using this relationship, it is easy to

fabricate TiN i thin ﬁlms with desired composition by adjusting the process parameters.

CHAPTER 2 LITERATURE REVIEW

2.1 TiNi Shape Memory Alloys

2.1.1 Shape Memory Properties

TiN i alloys are ordered intermetallic compounds based on the equiatomic
composition [6]. They experience a crystallographically reversible martensitic phase
transformation causing the Shape Memory Effect (SME). The SME can be described with
reference to the cooling and heating curves in Figure 2-1 [7]. There is no change in the
shape of a specimen cooled from above austenite ﬁnish temperature (Af) to below
martensite ﬁnish temperature (Mf). When the specimen is deformed below My, it remains
so deformed until it is heated. The shape recovery begins at the austenite start temperature
(As) and is completed at the austenite ﬁnish temperature (Af). At the inﬂection point
between A, and Af, about 50% of original shape is recovered. Also, the application of
stress can induce the martensitic phase transformation within a certain temperature interval.
The large anelastic strain (~ 5%) associated with the stress-induced transformation is

recoverable upon unloading, which give rise to the superelastic effect (SE).

2.1.2 Phase Stability and Occurrence in TiNi Alloys

From the TiN i phase diagram, as shown in Figure 2-2 [8], stoichiometric TiN i
phase exists at high temperature, with a steep solvus boundary on the Ti-rich side, slight
solubility for nickel, and only one intermetallic phase Ni3Ti on the Ni-rich side. The
diagram has traditionally showed an eutectoid decomposition (T iN i —> TizNi + Ni3Ti) at
903 K. However, Wasilewski et al. [9] proposed that the TiNi may remain to be a stable
phase down to room temperature with a very narrow stoichiometric range, as indicated in
Figure 2-3 [9]. Furthermore, Van 1.00 et al. [10], Nishida et al. [11] and Kim et al. [12]
reported the existence of metastable phases other than Ni3Ti on the Ni-rich side, such as

Ni3Ti2 and Ni14Ti11. Also, Saburi er al. [13] proposed that it might be more appropriate to

3

4
express Ni14Ti11 as a Ni4Ti3 phase based on the crystal structure. The equilibrium phase

diagram on the Ni-rich side for the low temperature range, however, is not well established
and controversy on the composition and structure of those phases exists. Most researchers
accept that the microstructures of TiN i alloys are primarily single phase, with small
amounts of other phases distributed throughout the matrix. Since the precipitate phases are
generally isomorphous with TiN i, they produce coherency strains which profoundly affect
transformation characteristics, as well as strength and ductility.

Melton [6] also pointed out that the presence of oxygen affects the microstructure of
TiN i alloys. According to the Ni-Ti-O phase diagram, Figure 2-4 [14], oxygen decreases
the stoichiometric range of the TiNi phase and can terminate compositions within a three
phase region consisting of TiN i, Ti4Nisz and Ni3Ti. Thus, the N i3Ti phase can be
present in a Ti-rich alloy and TizNi is sometimes found in Ni-rich alloys. Besides, the
oxide Ti4Ni20x is isostructural with the intermetallic TizN i, which has a lattice parameter
increasing slightly from 1.132 nm (TizNi) to 1.134 nm (Ti4Ni20x). This makes the

identiﬁcation of unique phases very difﬁcult.

 

S

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100% Martensite

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is
\

, 100% Austenite

   

Electrical Resistance

\\\<\\\\\\\\

S

 

//

s

Temperature

cool
above A: below M: below Nb below M: deformed

[ 1—>r 1—>A

SPECIMEN

 

 

 

 

 

 

I

original undeformed
shape

Figure 2-1 The shape memory effect is described with reference to a plot of
electrical resistance vs. temperature from which the characteristic transformation
temperature Mt, Mr, A: and A: are determined [7].

Weight. Percent Nickel

 

 

 

 

 

 

 

 

 

 

 

 

to So so to so so _ roo
' ' ' ‘ t
s
i
rseo°c
.. .1o°c
5:: 18°C
5 TiNi n
f; z
. o ‘0 i:
E
J.“
t-
. . t
? reef; g
600 ‘. . , . . , , , .
o 10 20 so 40 so so 70 so so 100
Ti Atomic Percent. Nickel Ni

Figure 2-2 Equilibrium Phase Diagram for TiNi Alloys [8]

 

 

 

 

 

 

 

 

 

I400
1200- / ”' \ \\ \\
w / >——-
O
I 1025 :20 /
I000~ " B
o B+TiNi3
0 -
3 800-
2
g 625 3 20°
‘” soc- ,’ . .
o. . . ~NI TI )
+ I ( 3 2
”0" , p+"Na,r.,
Ti Ni+mortensrte
z _ - ___________
o------‘:"‘7Z t .
45 50 55

Ni,otomic percent

Figure 2-3 Proposed equilibrium diagram of the Ti-Ni binary system in the
v1crmty of the equratomic composition at low temperature [9].

 

 

   
  

 
   

   
  
 
  

 

 

n o 1200 K
t 7
OTIINNOI
a-r. mm, n20,
OTINI,
a- n . rumor .Nim

B-Ti Ti,Ni m; TiNi, NiITi)

TI Ni

Figure 2-4 Isothermal Cross Section Through the Ti-Ni-O Phase Diagram at 1200 K
[14]

8
2.1.3 The Effects of Alloy Composition on The Transformation Temperature

The martensite start temperature (M,) depends strongly on composition, particularly
on the Ni-rich side, as shown in Figure 2-5 [6]. The M, temperature changes over 100
K/at% on the Ni-rich side, whereas the change is only about 20 K/at% on the opposite
side. On the Ti-rich side, the solubility range is almost independent of temperature.
Therefore, Ti-rich alloys show less sensitivity primarily as a result of the formation of a Ti-
1ich precipitate (TizNi), leaving the matrix composition essentially TiN i.

The strong dependence of M, on composition on the N i—rich side is a consequence
of higher Ni solubility at temperatures above around 773 K. Dissolved nickel may then
precipitate out upon aging for longer times at lower temperatures. The ﬁrst factor affecting
the M, is that the precipitation of Ni-rich phases leads to a more Ti-rich matrix composition,
with higher M,. The second factor is that the internal stress ﬁelds associated with semi-
coherent precipitates, such as Ni4Ti3, have a strong effect on suppressing the M:
temperature and constraining the further development of martensite [15]. Similar results
have been found when dislocations are introduced by cold work or transformation cycling
[1 6].

In addition, substituting 3d transition metals (such as Fe, Co, or Cr), and Cu, for
Ni has a great effect on the phase transformation behavior of TiN i alloys [2]. In these
cases, the effects are due to the incorporation of these elements into the lattice structure
itself and transformation temperatures are shifted. Also, the existence of interstitial
impurities like C, O, N, or H, which either form compound phases or exist inside the
lattice, will decrease the transformation temperature dramatically.

- Therefore, to maintain the best control over the transformation temperature, precise
composition control and avoidance of contamination by O or N are required when
fabricating the TiN i alloys. Depending upon the desired M,, the necessary composition

control is to between one tenth and one hundredth of an atomic percent.

 

 

 

 

°C
200 w I 7 I "1 n
O
150 - _ 0 . —
O
E 100 —- ' A _
a i A t ‘
8. t o. A
ii 50— o ——
E; l II {.A I
.8 AI;
<6 0 :‘A
E ‘ l* g
i “ '
E-' _50 l‘ ° Wang et al. A .
AiHarrison et al.
OlHanlon et al. A
400 _ .zPurdy et al. ‘ _j
. t
—150 I l J .1 l r

 

47 48- 49 50 51 52 53
Nickel atomic (%1

Figure 2—5 The Dependence of the Transformation Temperature (M1) on Nickel
Content in TiNi Alloys [6]

10
2.2 Mechanical & Metallurgical Studies of TiNi Thin Films

TiN i alloys not only possess shape memory and superelasticity effects, but also
provide high ductility, superior corrosion resistance, fatigue resistance, and high damping
capacity. Consequently, they have been used as high force actuators [l7], tubing
connectors [18], and orthodontic hardware [19]. In addition to development of
microactuators in microelectromechanical systems, TiN i thin ﬁlms have attracted attention
for potential application as a superelastic surface coatings to improve fatigue or erosion
resistance [3, 4].

TiN i thin ﬁlms with near-equiatomic overall compositions have been fabricated by
various methods, and their properties have been studied [20-32]. Thin ﬁlms formed by
vapor condensation in binary systems were predicted to be amorphous. Solid state
immisciblity and a large difference (greater than 10%) in the atomic radii of the constituents
were the most important ingredients for amorphous phase formation [20]. However, a
structural difference nrle has been proposed for amorphization by ion beam mixing.
According to this rule, the structural difference between the constituent elements of the
binary systems, rather than the radii or electronegativity difference, dictates amorphous
phase formation [21].

. High energy ion irradiation has been employed to mix bilayered nickel and titanium
deposited by electron beam evaporation. Rai er al. [22] reported that an amorphous TiN i
layer was formed after ion-mixin g. Ion beam sputtering and ion beam enhanced deposition
were other methods to make TiN i thin ﬁlms [4, 23]. The common advantage for electron
beam and ion beam processes is low process pressure. They can be operated at low
background pressure (~ 104 Pa) which results in less inclusion of gas molecules in the
sputtered ﬁlms and less scattering of the sputtered particles during transit. The
disadvantage is that it is difﬁcult to obtain a reproducible compositional results even at the

same process conditions.

l 1

Periodic multilayers of nickel and titanium deposited by dual magnetron sputtering
were studied by Clemens [24]. The structures of sputter-deposited multilayer TiN i ﬁlms
are strong functions of sputtering pressure. Low sputter pressure (~ 0.27 Pa) leads to
relatively high quality, textured crystalline layers. High sputter pressure (~ 1.36 Pa) results
in random orientation and poor layering, caused by island nucleation and growth. This
effect is due to that increase in pressure causing lower impact energy of the deposited
species, which results in lower atomic mobility. Amorphous TiN i thin ﬁlms with
compositions of Ni3oTi7o, Ni55Ti44 and Ni5gTi32 were also prepared by planar magnetron
sputtering [12]. The crystallization behavior of these thin ﬁlms were srn'veycd. For the
near—equiatomic thin ﬁlm (i.e. Ni56Ti44), complicated crystallization sequences and many
different precipitate phases were observed.

A triode magnetron sputtering method was adopted by Chang [25, 26] to fabricate
single layer (SL) of Ti(NiCu) and periodic multilayered (PML) of Ti(NiCu) and Ti thin
ﬁlms. Their compositions, crystallographies, interface coherency and corresponding
transformation behavior were studied. In compositional results, a severe Ti depletion,
relative to the sputter-target composition, was found in the SL ﬁlms. The reason causing
this phenomenon was believed that the growing ﬁlms were bombarded by energetic
species, resulting in preferential resputtering. According to the published sputtering yield
data of elemental Ni and Ti [40], Ni shows higher sputtering yield than Ti at the same
bombarding energy. Thus, a Ni depletion, produced by preferential resputtering. is
expected in the growing ﬁlm. However, the Ti loss found in the SL ﬁlms in Chang's
studies indicated that the sputtering properties of Ni and Ti in amorphous TiNi alloys may
be different from pure elements and crystallized TiNi target alloys.

Ion beam assisted deposition (IBAD) experiments were thus designed by Chang
[26] to gain better understanding of the sputtering behavior of Ti and Ni in amorphous TiNi
alloys. The Ti depletion is plotted as a function of ion-to-atom (I/A) arrival ratio, shown in

Figure 2-6. The results showed that a general trend of Ti depletion in the IBAD processed

12
ﬁlms which increased with increasing of VA ratio. Also, a dramatic increase of the

sputtering yield ratio Y = YT'JYN, is found at low assist-beam energy. In the low energy
regime (~ 50 eV), the sputtering rate is extremely sensitive to the variation of ion energies
(E;) and energy of sputtering threshold (E,;.) [27] (threshold energy, the minimum
requirement energy for sputtering). Since Ni has a higher threshold energy than Ti, as
shown in Table 2-1 [28], the increase of sputtering yield ratio can explain high Ti depletion
rate for the 50 eV assisted-beam energy. Chang [26] concluded that the relation of the
sputtering yields of Ni and Ti in binary amorphous TiN i alloys can be expressed as YT, >
YN; in low and intermediate energy regimes (< 500 eV), and the sputtering yield ratio of Ti
and Ni in amorphous TiN i alloys increased ﬁom 1.75 at 500 eV to about 9 at 50 eV of Ar

 

 

 

bombarding ions.
Table 2-1 Threshold Energies (eV) [28]
Ne Ar Kr Xe Hg__

Al 13 13 15 18 18
Ti 22 2 0 17 18 25
V 21 23 25 28 25
Cr 22 22 18 20 23
Fe 22 20 25 23 25
Co 20 25 22 22

Ni 23 2 1 25 20
Cu 17 17 16 15 20
Nb 27 25 26 32
Mo 24 24 28 27 32
Ag 12 15 15 17

W 35 33 30 30 30
Pt 27 25 22 22 25
Au 20 20 2O 18

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

13
22.2 9 (500 9V) 19
2w 22.2 a (wow) 9
° ! i l 1 I .
C i i E 0 50 eV 2
'0-02 L'" """ i """""" i """""" i ------- I 100 eV ------ -:
: .f . i i o 500 eV 1
‘0-04 I """" : “""i': """""" i """"""" J; ----------- 1 ---------- 1
P 5 ‘I E 5 5 .
A P ' t t g «-
.\° -o.06 :- -------- -‘ ----------- i -------- . - . ------- : ----- :
u 1 I 9 I
c : s E a a s :
i: ’0-03 .7 """"" i """"" i """""" 1 """"""" 1 ----------- T --------- 1
<1 - : s a a a -
-o.1 :- ---------- ;~ --------- 4 ----------- i ----------- i ----------- ~5- --------- —.
~0.12 .- """"" E‘ """""" i """"" ‘2 """""" ‘2 """""" ‘i """"" 1
. r r a = a r
-0.14
0 0.1 0.2 0.3 0.4 0.5 0.6

Ion/Atom Arrival Ratio

Figure 2-6 The change of Ti concentration of the [BAD processed TiNi ﬁlms as a
function of HA ratio, relative to the expected composrtron of the ﬁlm 1n the absence
of an assist-beam. Data are shown for 50, 100, and 500 eV beam [26].

14

Since 1988, Busch, Johnson and his coworkers have conducted research on shape
memory properties and applications in TiNi thin ﬁlms [3, 29, 30]. TiNi thin ﬁlms were
deposited from an alloy target by dc magnetron sputtering. Sputtering targets were
prepared from different compositions of TiNi alloys. One target with composition Ti -
50.1 at% Ni had a transition temperature of 303 K. The other target with composition Ti -
50.2 at % Ni had a transition temperature slightly above 373 K. The sputtering condition
were: PA, = 0.1 Pa, I = 0.5 A, V = 450 V, source/target distance = 57 mm. After
deposition, the resulting ﬁlms were amorphous. Their crystallization temperatures were
determined by Differential Scanning Calorimetry (DSC) and showed that the crystallization
exothenn was located at about 753 K. Knowing the crystallization temperature, samples
were vacuum encapsulated in quartz ampoules with a Tr sponge getter, annealed for 30 rrrin
at 823 K to ensure crystallization complete and air cooled. DSC was also used to identify
and evaluate their phase transformations temperature behavior. It was found that the
transformation temperatures of thin ﬁlms were about 100 K lower than their target alloys.
This phenomena was described as being the contamination of oxygen since oxygen can
change the stability of TiN i phase. The composition of the target materials and the
respective ﬁlms were measured with the Flame Atomic Absorption method. Questions
exist on the results: 1) the measured target compositions were signiﬁcantly different from
the nominal compositions; 2) the target alloys were richer in Ni than their thin ﬁlm
counterparts, but their transformation temperatures were higher. This differs ﬁom the
general rule stating that as Ni concentration increases, M, decreases. The 1eaSons behind
the results were not clariﬁed.

Even though low transformation temperatures existed, the classical shape memory
behavior was successfully demonstrated on TiN i thin ﬁlms prepared by sputter deposition
and a crystallization process. Two devices were designed by Johnson [3] using the SME
of TiNi thin ﬁlms: one is a microvalve and the other is nricroactuator, as shown in Figure

2-7 and Figure 2-8.

 
  
 
  
 
  
 
 

Brass Cover_
Plate

Formed in End

Silicon Valve —
Substrate

Outlet J

_ __Nylon Turned
' Cover

_ Membrane Bias
Pressure

Kapton Insulation

L Outlet

Figure 2-7 The Valve Design by Using TiNi Thin Films [3]

Semiconductor Laser

Film Heated
to 150 °C

Laser Heats One of the
Two Actuators

Unheated Film '
Plastically Dcfomls 11335353!“

(APPI'OX- 3%) Annealed Shape

   

"1" or "0" Determined and Set
by Which Actuator is Used

Figure 2-8 The Mirror Actuator Concept by Using TiN i Thin Films in [3]

16

Ishida et al. [31] used the rf magnetron sputtering method to fabricate TiN i thin
ﬁlms under different sputtering conditions. The advantages of rf magnetron sputtering
over dc conﬁguration will be discussed in Section 2.4.4. His compositional results were
measured by electron-probe microanalyzer (EPMA) and shown in Table 2-2. Note that at
low pressure process condition (Run 5), the composition was Ti depletion, too. This can
strengthen the results in Chang's study [26]: at low Ar gas pressure, the resputtering of
TiNi amorphous thin ﬁlm resulted in the Ti loss in the thin ﬁlms. Also, The mechanical
properties of the annealed ﬁlms were evaluated. They showed that the ﬁlms formed at a
high Ar gas pressure (13.3 Pa) exhibited poor mechanical properties due to the porous
structure formed during sputtering. The ﬁlms prepared at a low Ar gas pressure (0.67 Pa)
exhibited a good shape memory behavior comparable with that of bulk TiN i alloys.

Table 2-2 Sputtering Conditions and Film Compositions in [31]

- (Distance between target-substrate was 82 mm, and arget composition is Ti-50 at% Ni.)

 

 

 

 

Run Ar gas R.F Substrate Film Composition (at%)
Pressure Power temperature
Pa W °K Ti Ni
1 13.3 400 523 50.2 49.8
2 6.7 600 573 50.5 49.5
3 6.7 400 523 50.6 49.4
4 6.7 200 523 51.6 48.4
5 0.67 400 523 48.6 51.4

 

 

 

 

 

 

17
Walker et al. [32] deve10ped and demonstrated the deposition of thin ﬁlm shape

memory alloys and an associated dry etch, surface micromachining release process. The
deposition and patterning procedure may be used for the materials and processing
technology of silicon-based microelectronics. In addition, the activation and control of
such structrn'es is possible using currents and voltages compatible with standard IC
electronic circuitry. However, the lack of detailed process flow makes the results suspect
because the ﬁlm should be amorphous in these experiments.

In summary, the research on TiN i thin ﬁlms is drawing much attention and broad
applications are developing. However, methods for fabricating thin ﬁlms are critical since
the transformation temperatures of TiN i alloys are extremely sensitive to Ti concentration.
How to make a thin ﬁlm with a precise composition control is the major concern in the
fabrication processes. Sputtering is believed the best method to fabricate alloy thin ﬁlms
with close composition control. Nevertheless, the different process conditions for
deposition will result in different composition results. The present work is designed to

address on the relationship between process parameters and ﬁlm compositions.

2.3 Deposition by Sputtering
Thin ﬁlms can be fabricated by many methods, including thermal evaporation,
sputtering, ion beam deposition and chemical vapor deposition. Among these methods,
sputtering deposition has become one of the most versatile techniques in thin ﬁlm
technology for preparing thin solids ﬁlms of almost any material because of the following
advantages [33]:
(1) Sputtering can be accomplished from large-area targets, simplifying the problem of
depositing ﬁlms with uniform thickness over large substrates.
(2) Film thickness control is relatively easily achieved by selecting a constant set of
operating conditions and then adjusting the deposition time to reach the desired ﬁlm
thickness.

18
(3) The alloy composition of sputter-deposited ﬁlms can be more tightly controlled than
that of evaporated ﬁlms. This is especially important for alloys and compounds.
(4) Many important ﬁlm properties, including stress and adhesion, can be controlled by
altering process conditions, such as power and pressure.
(5) Once the process conditions are ﬁxed, the same quality of thin ﬁlms can be obtained
repeatedly.

2.4 Principles of Sputtering

In general, sputtering is a term used to describe the mechanism by which atoms are
dislodged from the surface of a material (target) by collision with energetic particles. These
ejected atoms will condense on any solid substrate close to the target. The sputtering
process consists of four steps: (a) ions are generated and directed at the target; (b) the ions
sputter target atoms; (0) the ejected (sputtered) atoms are transported to the substrate,
where; ((1) they condense and form a thin ﬁlm. It is important to note that the substrate may
also be subject to energetic particle bombardment, and atoms may be sputtered from the
substrate as well. It is in this context that differential sputtering rates complicate

composition control.

2.4.1 Properties of Glow Discharge

The energetic particles used to strike target materials are generated by glow-
discharge. A Glow-discharge is a self-sustaining type of plasma (a partially ionized gas
containing an equal number of positive and negative charges, as well as a number of non-
ionized neutral molecules) [33, 34, 35]. Figure 2-9 [34] is an illustration of the appearance
of glow discharge for a simple dc diode type system. It consists of a glass tube which is
evacuated and then ﬁlled with a gas at low pressure. Within the tube are two electrodes

between which a dc potential difference is applied.

l9

 

 

 

 

 

 

 

 

 

GLASS ENVELOPE
CATHODE MODE
H |%l/\ \\\\\ m
CROOKES NEGATIVE FARADAY POSITIVE \E ANODE
DARK I GLOW DARK COLUMN :
SPACE SPACE : SPACE
I
II (a) i
l
I
I
| .
I
n+ (—) I 5
. I
n-(...) "v ' ........................................................ ..:
I
t o \
i.
ll '
I
I
(b) g
i
ll /:
I
VOLTAGE I i
E
I
II =
i
(C)
Figure 2-9 [34]

(a) Structure of a Glow Discharge in a DC. Diode System
(b) Charged Particle Concentration in a Glow Discharge

(0) Voltage Variation in a DC. Diode Glow Discharge

20

If a free electron is introduced into the tube (mostly likely created from the
ionization of an argon atom by a passing cosmic ray), it will be accelerated by the electric
ﬁeld existing between the two electrodes. Then, it will pick up energy high enough to be
transferred to the orbital electrons of the Ar atoms during an inelastic collision, causing
their excitation or ionization. If the transferred energy is less than the ionization potential,
the orbital electron will be excited to a higher energy state for a short time. The electron
will then return to the ground state simultaneously emitting visible-light photons. Such
excitation is the source of light emission in glow discharges. If the energy transferred is
greater than the ionization potential, a second free electron (and positive ion) will be
created. Subsequently, both free electrons will become accelerated again, thus cascading
the number of free electrons creating a condition known as gas breakdown.

When the condition of gas breakdown is reached, electrons and ions are quickly
lost to each of the electrodes and to all other surfaces within the chamber. Such loss
processes include electron-ion recombination and an equivalent electron loss into the
external circuit at the anode. Therefore, the current will decay to zero and glow discharge
will extinguish unless there is a mechanism available for generating additional free electrons
for sustaining the current flow. When a sufﬁcient number of electrons are available to
maintain the discharge, the discharge is said to be self-sustained.

The self-sustaining discharge has a particular structure, as shown in Figure 2-9a.
The most important region of the discharge is the Crookes dark space between the negative
glow and cathode. Any electrons near the cathode are rapidly accelerated away from it, due
to their relatively light mass. The much more massive ions are accelerated toward the
cathode, but much less quickly. Thus, on average, they spend more time than electrons
traversing the Crookes dark space, and at any instant their concentration in the dark space is
greater than that of electrons (The dark space can also be explained by the lower density of
electrons, which causes lower excitation of atoms). This has the net effect of greatly

increasing the electric ﬁeld immediately in front of the cathode. As a result, the electric

21
ﬁeld in the remainder of the discharge is rather low and uniform. The greatest part of the

voltage, between the anode and cathode, drops across the Crookes dark space so that
charged particles experience their largest acceleration in this region.

When positive ions from the negative glow region enter the Crookes dark space,
they experience the strong local electric ﬁeld and become accelerated toward the cathode.
They also have a high probability of exchanging their charge with neutral atoms in the dark
space. In doing so, they retain their momentum, while losing their charge. The formerly
neutral atom becomes an ion, and only at that moment does acceleration towards the
cathode begin. Virtually no ions reach the target with the full dark space energy. In
addition, this implies that the cathode is bombarded by energetic neutral atoms as well as
energetic ions, and thus both species produce sputtering events.

The source of electrons that sustains the discharge is emission by the cathode of
secondary electrons when struck by ions. When the electrons enter the dark space from the
cathode, they are accelerated by the dark space and result in ionization. Then, new
secondary electrons can be produced by ion bombardment of the cathode again. Thus,
there only needs to be a sufﬁcient number of ions produced in the dark space to produce
enough secondary electrons to sustain the glow discharge.

If the pressure becomes too low (< 1.3 Pa), there are insufﬁcient ionizing collisions
between the electrodes, and the glow discharge extinguishes. Therefore, for the sputtering
system to operate at pressures lower than 1.3 Pa, there should be another electron source or
the ionizing efﬁciency of the available electrons should be increased, such as can be
accomplished by the application of a magnetic ﬁeld.

2.4.2 Physics of Sputtering
When a solid surface (target) is bombarded by atoms, ions, or molecules, many
phenomena can occur, as shown in Figure 2-10 [34]. At very low energies (< 5 eV), ions

are likely to be reﬂected or neutralized in the process. At much higher energies (> 10 keV),

22
the impinging particles are most likely to be embedded in the target. This mechanism is the
basis of ion implantation. At energies between the two extremes, some of the impinging
ions may be responsible for the formation of an altered surface layer, i.e. structure
rearrangement in the target material. Other ions result in a series of collision between
atoms, leading to the ejection of one of these atoms from the surface into the gas phase
(sputtering).

G.K. Wehner [36], whose theoretical work established a scientiﬁc basis for
sputtering, often described sputtering as a game 'of three-dimensional billiards played with
atoms, as shown in Figure 2-11. The sputtering process is compared to the "break" in a
game of pool, in which the cue ball strikes the orderly arranged pack, scattering balls in all
directions, including those going back to the player. Using this analogy, it is possible to
realize how atoms may be ejected from a surface as the result of two binary collisions when
a surface is struck by a particle with a velocity normal to the surface.

The sputtering yield, deﬁned as the number of atoms ejected from the target surface
per incident ion, is the most important parameter for characterizing the sputtering process.
When the directions of sputtered atoms from the surface of polycrystalline materials are
measured for the case of normal incidence, it is found that the directions of the ejected
atoms leaving the surface essentially follow a cosine distribution. Evidently, in actual
sputtering events, more than two collisions are involved, and the energy delivered is so
randomly distributed that the effect of the incident momentum vector is lost. For the case
when the srn'face is bombarded by ions at an oblique angle, there is a higher probability that
the primary collision between the incident ion and the surface atom will lead to a sputtering
event. Furthermore, oblique incidence conﬁnes the action closer to the surface, and thus
sputtering is enhanced. Thus, the sputtering yield is a function of the incident angle with
respect to the normal of target surface, as shown in Figure 2-12 [37]. It follows the sec(0)
curve and attains its maximum at 0 of about 70°. This can be used to account for the

different sputtering behavior for differing target morphologies.

23

 

 

Incident Reﬂected Ions
Ion & Neutrals
Secondary
Electrons
Sputtered
Ams
Surface
Structural
Change Possible

Bombarding Ions
May be Implanted
Collision May Terminate or Result in The Ejection
Sequence: Within the Target of a Target Atom
’ (Sputtering)

Figure 2-10 Interactions of Ions with the Surface [34]

24

 

 

     

NEUTRALIZED

GAS MOLECULE
REBOUNDED FROM
THE TARGET SURFACE

BOMBARDING IONIZED

GAS MOLECULE

b )
TARGET ATOMS \
—-» INDICATE DIRECTION OF SPUTTERED ATOM
MOMENTUM TRANSFER

Figure 2-11 (a) Binary Collision Between Atom A and B, Followed by a Binary
Collision Between Atom B and C. (b) Collision Process Responsible for Sputtering
and Fast Neutral Generation [36]

25

 

 

 

 

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E ,0 fo'O . .2
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i o
0.5 f— k.
' a
i l i
O 30 60 90

ANGLE OF INCIDENCE (degrees)

Figure 2-12 Variation of Sputtering Yield with Angle of Incidence for 1 keV Ar Ion
Incident on Ag, Ta, Ti, and Al [37]

26
The sputtering yield also depends on a number of factors besides the direction of

incidence of the ions; it depends also on the target materials, the target composition and
structure, the mass of bombarding ions, and their energy. There is a minimum energy
threshold for sputtering every material, as shown in Table 2-1. In the energy range of
sputtering (10-5000 eV), the yield increases with ion energy and mass. Figure 2-13 [27]
shows the sputtering yields of nickel as a function of energy for various noble gas ions.
The sputtering yields of various materials in argon, at different energies, is given in Table

2—3 [36].

Table 2—3 Sputtering Yield for Metals in Argon (atoms/ion) [36]

 

 

Target 100 eV 300 eV 600 eV 1000 eV 2000 eV
Al 0.11 0.65 1.2 1.9 2.0
Au 0.32 1.65 2.8 3.6 5.6
Cu 0.5 1.6 2.3 3.2 4.3
Ni 0.28 0.95 1.5 2.1
Pt 0.2 0.75 1.6
Si 0.07 0.31 0.5 0.6 0.9
Ta 0.1 0.4 0.6 0.9
Ti 0.08 0.33 0.41 0.7
W 0.12 0.41 0.75

 

27

 

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Figure 2- 13 Sputtering Yield of Nickel as a Function of Ion Energy and Ion Mass
[27]

28
Several matters related to sputtering yield should be noted. First, although the

Sputtering yields of various materials are different, as a group they are much closer in value
to one another than the vapor pressure of comparable materials. This makes the deposition
of multilayer ﬁlms or multi-components ﬁlms much more controllable by Sputtering than by
evaporation. Second, since the bombarding ions are not monoenergetic in glow discharge
sputtering, it is not necessarily valid to use the Sputtering yield value for pure metals when
alloys and compounds are being sputtered. Third, since the mass differences between
components, the momentum and energy will be distributed differently to each constituent
resulting in different ejection probabilities for the atoms. Note that the sputtering yield ratio
between the alloy components doesn't equal to the deposited materials composition ratio
since there are many other factors will change the ﬁnal composition outcome. This will be
discussed in section 4.1. The tabulations of Sputtering yields, however, are useful for
obtaining rough indications of deposition of various materials.

For the deposition process, argon was chosen as the sputtering gas for its inert
properties and low cost. The pressure range of operation is set by the requirements of the
glow discharge (lower limit) and the scattering of sputtered atoms by sputtering gas (upper
limit). Generally Speaking, the higher the current at the cathode, the higher is the
deposition rate (since more ions are striking the cathode, and thus are causing more
Sputtering).

In fact, sputtering is a highly inefﬁcient process. About ~ 70% of the energy
consumed during the sputtering process is dissipated as heat in the target, and ~ 25% by
emission of secondary electrons and photons by the target. The target heating can raise
target temperatures to levels capable of damaging the target, associated vacuum component
and the backing electrode. Cooling of the target with water is typically used to maintain

low target temperatures.

29
2.4.3 Sputter Deposition Film Growth

Upon being ejected from the target surface, sputtered atoms have energies about 10

to 40 eV. In general, the target-to-substrate spacings are 5-10 cm. The mean free path, A,

of a sputter gas atom under typical sputter pressure is 1 cm for Ar atoms at 0.6 Pa, less

than 5-10 cm. (The mean ﬁee path is calculated from equation it = ——-1—, where n is

J2Mld2n
the gas concentration, d, and d, are the atomic diameters of colliding atoms, one of them
being the background gas atom's [34]). Thus, it is likely that the sputtered atoms will
suffer collisions with the sputter gas atoms before reaching the substrate. The sputtered
atoms may therefore arrive at the substrate with reduced energy and increased energy
spread; be backscattered to the target or the chamber walls; or lose enough energy so that
they are transported by diffusion in the same manner as neutral sputter gas atoms [34, 35].
As these events occur during the transport of sputtered atoms to the substrate, sputtering
gas pressure can alter the energy of arriving matter and thus affect the properties of the
deposited ﬁlms.

The formation and growth of thin ﬁlms on the substrate follows mechanisms of
nucleation and growth, as shown in Figure 2-14.[38]. Atoms transfer the normal
component of their velocity to the substrate and attach to the surface. After becoming
attached to the surface, they are refened to as "adatoms". Adatorns may continue to move
along the surface due to the kinetic energy associated with their initial lateral velocity, or by
thermal activation from the surface. In some cases, the adatoms may re-evaporate from the
surface into the vapor phase. AS the adatoms migrate along the surface, they may interact
with each other to form nuclei. In the formation process, the ﬁlm passes through stages of
nucleation —> growth —) island formation —) island coalescence -) continuous ﬁlm

formation.

30

(a) Single Atom Arrives (e) Growth

 

Islands Growing

 

 

 

 

 

 

Substrate
(b) Migration Re-evaporation (1) Island Shape
. Cross-Section

 

 

 

 

 

 

 

(c) Collision & Combination of
Single Atoms

/ I e I e I‘
- Go
‘I I I I
I
obeﬁeﬂeﬁoﬂ
.Jlfifififit.

' .
CD ~iPJ-1'J-Z'J-Z'ﬁ'tJ-Bhn
mini-rarest.

(g) Coalescence

 

 

 

05%
S
I
‘-
I
S
I
‘-
I
‘-
o'

 

 

 

 

 

 

 

(d) Nucleation (h) Continuity

"Island" of Atoms
%
%

 

 

 

 

 

Figure 2-14 Formation of Thin Film [38]

3 1
The substrate on which we desire to deposit a ﬁlm is important in determining the

quality of the ﬁlms. Some factors inﬂuencing quality are the substrate's temperature,
chemical nature, crystallographic structure and condition of the surface. Also, the substrate
is subjected to impingement by many species during this process, as shown in Figure 2-15
[34]. The most important species are: 1) fast neutral sputter gas atoms (Ar), which retain
signiﬁcant energy after having struck and recoiled ﬁom the cathode. As they impinge upon
the substrate, some may embed themselves in the growing ﬁlm. 2) Negative ions are
formed near the cathode surface by the reaction of secondary electrons. The latter can
acquire substantial energy from Crookes dark space electric ﬁeld, and thus also Strike the
substrate with appreciable energy. 3) High energy electrons, which are accelerated across
the Crookes dark Space, and then impinge on the substrate, and: 4) Low energy neutral
sputter gas atoms (Ar), which strike the substrate at a high ﬂux. However, the Sticking
coefﬁcient of low energy neutral Ar atoms is negligibly small, and thus the only
incorporation of Ar atoms into the growing ﬁlm arises from fast neutral embedment.
Finally, 5), contaminants present in the form of residual gases in the chamber can be

incorporated into the growing ﬁlm.

32

Argon:
Thermal Neutrals
Fast-Target
Fas+t+—Charge Exchange

Complexes

 
  
   

 

Contaminants

    

 

5 Photons ». , if; EleCtrons,
Sputtered / . Negatrve _,_ é
Atoms ,. , v- "

Ions

 

 

 

 

 

 

 

Figure 2-15 Particles Bombarding the Substrate in Sputter-Deposition [34]

33
2.4.4 Magnetron Sputtering

In general, the diode conﬁguration is the simplest glow-discharge based sputtering
system and possesses severe limitations that make it unsuitable for most practical
applications. As a consequence, radio-frequency (rf) sputtering, and magnetron sputtering,
are most often used to deposit thin ﬁlms. The rf sputtering method is used for the
deposition of insulator materials since the glow-discharge cannot be maintained with a dc
voltage if the electrodes are covered with insulating layers. Also, it can be used to deposit
conductive materials. The most common type of rf magnetron conﬁguration uses the
chamber and other grounded ﬁxtures for the second rf electrode. With this single-ended
conﬁguration, the area of the Sputtered target surface is usually small compared with that of
the ground electrode, and only the target electrode has a sufﬁciently large bias voltage for
spuuering. If the length of the magnetron is less than 30 cm, rf sputtering is as easy to
implement as dc sputtering.

In both dc and rf diode sputtering, most secondary electrons emitted from the target
do not cause ionization events with Ar atoms. They end up being collected by the anode or
at substrates where they may result in unwanted heating. The application of a magnetron
increases the ionization efﬁciency of elements by utilizing magnetic ﬁelds to help conﬁne
the electrons near the target surface and lengthen their paths by conﬁning them to Spiral
orbits.

To describe the principle of the magnetron conﬁguration, ﬁrst consider the motion
of an electron, mass m,, and initial velocity v, that is perpendicular to a uniform magnetic
ﬁeld, as shown in Figure 2—16a [34]. The electron experiences a force, F = qv x B ,
which is perpendicular to both its velocity and the direction of B. If no other forces are
exerted on this electron, it will continuously be acted upon by the force F, and this will

induce a circular motion with radius:

r=(m‘v/eB)

34
Comparing electrons and Ar ions under equivalent conditions of B and particle energy, the

equation for r shows that the Ar ions have motion radii which are about 300 times larger
than the electrons' due to the signiﬁcant difference between their masses. This illustrates

that the direction of electron motion in magnetron discharges is strongly inﬂuenced by the
magnetic ﬁeld, while the magnetic ﬁeld does not Signiﬁcantly change the direction of Ar

ions as they cross the dark space.

Let there be a magnetic ﬁeld, B, parallel to the swim and an electric-ﬁeld§
perpendicular to the surface, as shown in Figure 2-16b. Let the electric ﬁeld E decrease
linearly across the dark space of thickness L with 50 at the surface of target. If y is the
vertical distance above the target and the target surface is y = 0, then the magnitude of the

electric ﬁeld is given by:
§= E, (1- y / L)

Assume the electron is initially at rest. Now the electron will be accelerated away
from the target in a direction perpendicular to the surface by the electric ﬁeld, but
simultaneously it will experience a force due to the magnetic ﬁeld, F = qv x B. The
electron velocity will therefore be altered from a direction perpendicular to the surface. If
the magnetic ﬁeld is strong enough to deﬂect the electron velocity so that it begins to return
to the target surface before it leaves the magnetic ﬁeld, its motion will be roughly cycloid as
shown in Figure 2-16b. That is, as it approaches the target sm'face after having been
deﬂected, the electron will be decelerated by the electric ﬁeld, and eventually will come to
rest (momentarily) at the target surface. At this point, the next period of the cycloid motion
will be initiated, as it is once again repelled by the electric ﬁeld. The maximum excursion

of the electron from the target during such motion is given by [39]

“2
1 2m.
y... = - (V - V,)
B e

 

where V, is the potential of target and V is the potential in the dark space at ym.

35
The net result is that secondary electrons are trapped near the target surface by the

combination of the magnetic and electric ﬁelds. Even though not every electron produces
an ionization event, many more ions are generated than if the secondary electron escaped
the discharge without the application of the magnetron.

In a planar magnetron [40], the target surface is planar, and the B-ﬁeld is created by
magnets behind the target, as shown in Figure 2-17 [34]. In examining these illustrations,
it must be noted that the magnetic ﬁeld lies parallel to the target surface half-way between
the magnetron poles. In this region, the E x B ﬁeld closes on itself. Thus, a continuous
path for the "hopping" electrons is established. This region is called the "racetrack" and
sputtering is very rapid in the racetrack portion of target. This will result in localized
erosion of target, leading to loss of absolute planarity. In addition, since magnetrons are
much more effective in causing secondary electrons to undergo ionization collisions, much

higher currents occur for comparable voltage levels.

36

 

 

 

 

 

 

 

 

 

 

i=0
V O V
B
.................. ,
IIIIIIIIIIIIIIIIII[ m,v =B¢VIIIIIIIIIIIIIIIIIII
IIIIIIIIIIIIIIIIII r IIIIIIIIIIIIIIIIII
IIIIIIIIIIIIIIIIII ___m,v IIIIIIIIIIIIIIIIII
IIIIIIIIIIIIIIIIII Be IIIIIIIZIIIIIIIIII
(a)
_ Y
E €00"?
Y
(b)

Figure 2- 16 Motion of an Electron Ejected from a Surface with Velocity v into a
Region of Magnetic Field B Parallel to the Surface [34]:

(a) with No Electric Field

(b) with a Linearly Decreasing Field

37

  
  
 

Magnet Poles

Hopping Electrons

 

Magnetic Field Lines

 

Target Material

 

 

 

 

Figure 2- 17 Schematic Drawing of Planar Magnetron Target and its Cross Section [34]

38
2.4.5. Deposition of Alloys Film

When an alloy surface is subjected to energetic particle bombardment, the different
elemental components are not sputtered in proportion to their concentration in the alloy [34,
35, 41]. This phenomena is called preferential Sputtering. The preferential removal of one
component from the alloy surface results in the formation of an altered layer, which is a
near-surface region with a composition different from the bulk alloy and depleted in the
Species experiencing the higher sputtering yield [41].

If a virgin target were used for sputtering, the layer of metal oxide will be removed.
Dming this period, the removal of oxide will be accompanied by a change (usually
decrease) of discharge current because of differing secondary electron coefﬁcients for metal
and oxide. After time, the nominal composition surface is exposed. Due to the preferential
ejection of the high sputtering yield component, the sputtered ﬂux leaving the surface
would be different from the composition ratio of target. But, barring rapid diffusion in the
target, this situation cannot last for long because the altered layer will become enriched with
the low sputtering yield component. Ultimately, a steady-state can be set up, whereby
elements are sputtered from the target surface in the same ratio as are present in the bulk
target.

Note that the depletion of one component on the target surface will produce a
concentration gradient of the components, and this will encourage diffusion. If diffusion
happened during the process, steady-state can't be attained. At sufﬁciently low
temperatures, the thermal diffusion can be minimized. So, it is very important to keep the
target cold enough during the sputtering process.

Transport of the Sputtered material from target to substrate is not merely a Straight
line travel. The collision between sputtered atoms and sputtering gas atoms will take place.
Especially at high pressure, transport becomes more like a diffusion process. Also, as the
mean free path of Sputtered atoms decreases, more of them will land on other places inside

the chamber and proportionately less on the substrate. Because of different mean free paths

39
for different components, which depends on atomic radius, the ratio on the substrate is not

expected to be the same as the ﬂux leaving the target surface.

The next stage is the condensation of the vapor atoms on the substrate. The
condensation coefﬁcient is the proportion of the atoms that stays on the substrate without
evaporation. This is determined by the arrival rate, temperature, and the bonding energy
between substrate and adatorrrs. Usually, the condensation coefﬁcient is different for each
constituent causing an effective change in the composition of the depositing ﬁlms.

In the glow discharge environment, the growing ﬁlm will be bombarded by ions
and to some extent there will be re-sputtering of the growing ﬁlms. Since sputtering
increases with increasing ion energy, such resputtering will be particularly effective for
high energy sputtering. In addition, the ﬁlms will be depleted of the component with the
higher sputtering yield due to the resputtering.

Considering all of these various effects together, the composition of a sputter
deposited alloy ﬁlm can be different from that of the target. However, the composition
change should be reproducible for a given set of process conditions. So by adjusting the

process parameters, thin ﬁlms with different compositions can be easily obtained

2.5 Composition Analysis

In this study, the energy dispersive X-rayspectrometry (EDS) method was adopted
to measure the composition. Because the necessary composition control in the deposition
of TiN i thin ﬁlms is to between one tenth and one hundredth of an atomic percent, the
accuracy of composition measurement thus becomes very important. In this section, the
principle of EDS analysis and the so called ZAF correction (See section 2.5.3) are
thoroughly discussed. In order to obtain an accurate compositional result of thin ﬁlm

without interfering by the substrate, there is a criterion for the thickness, about 3 um.

2.5.1 Principle of EDS Analysis

The scanning electron microscope-electron probe microanalyzer has been used to
determine the compositions of thin ﬁlms. As atoms are excited by electron beams,
characteristic X-rays are emitted when an electron jumps from one shell to another of a
lower energy state. By measurement of the energy or wavelength and number of X-rays
emitted by emission or ﬂuorescence spectrometry, compositional information in terms of
the atomic species present can be obtained Presently, the emitted X-rays are measured
with an energy dispersive X-ray spectrometer (EDS) [42, 43, 44].

The relationship between the intensity of an X-ray line and the concentration of the
element concerned is very complex. Information about the system's geometry, the
behavior of the specimen and its constituent elements are necessary. Thus, complex
mathematics and statistics must be applied to transform the X-ray raw data into the
quantitative results.

As Shown in Figure 2-18 [45], the electrons' distribution and the activated volume
in thick and thin specimens are different. With thin specimens, X—rays can usually be
regarded as having come from an area the Size of the incident electron beam diameter. With
thick, bulk specimens, determine this area of origin is much more complicated because
electrons interact with a larger volume of the Specimen. This interaction volume depends
on several variables, such as energy of incident electrons, materials, and tilt angle. Thus,
bulk samples will produce much larger error in quantitative work than thin specimens.

The most important technique for analyzing quantitative electron-excited results is
known as a ZAF calculation. ZAF is an acronym for the three effects: atomic number (Z),
absorption (A) and ﬂuorescence (F). Noting that the Z correction is a combination Of the
'stopping power' and the backscatterin g correction, both of which are primarily dependent
on atomic number. After ZAF corrections, the specimen/Standard intensity ratio can be

converted into concentration.

41

Bulk Sample l l l

    
 

Detection

 

 

Fluorescence

1

Thin Foil Sample

Detection

\‘AA\\\\\\\\‘;/ \\\\\\\\\\3\

 

 

 

Z

Figure 2-18 Schematic Distribution of Electrons and Activated Volume in a
Thick and Thin Specimen [45]

42

2.5.2 Characteristic X-ray Intensity

The number of ionization (dn) per increment of electron path length (dx) on a pure
element is

dn = Q I: 9% dx

where Q], is the ionization cross-section of the shell concerned (usually the Shell K), N is
Avogadro's number (6.02 x 10” ), p is the density, and A is the atomic weight of the
element concemed. Assuming that incident electrons lose energy continuously along their

path, there is a relationship between path length, x, and energy, E. By integration, the total

number of ionizations produced along the trajectory of an electron with initial energy E, is :

50 dix
n=—lQ,——ds
5, A (II?

where E, is the threshold excitation energy. The minus Sign considers the decrease in E
with x that makes dx/dE negative.
Deﬁning the "stopping power" (S) of the bombarded material as —dE / d ( px ),

then
N E. Q
n = —. l —"a
A E, S
multiply n by R (a factor taking into account the loss of ionization caused by electron

backscattering) and 0),, (ﬂuorescence yield, the fraction of ionization of K Shell that result in
characteristic X-ray emission) gives the characteristic X-ray intensity per incident electron
on pure material
E
N 0.) [C Q It 0 R
1,, = — l —dE
A BC S

where to], is a constant, Q], is a function of E0 and E, only.

43
2.5.3 ZAF Correction
If compounds were considered, the term Np/A (the total number of atoms per unit
volume in pure element) must be replaced by CNp/A (the number of atoms of the element
concerned per unit volume in the compound), where C is the concentration of that material.
The ratio of characteristic X-ray intensity (I) from the specimen to that emitted by a

standard (Io) containing a concentration Co of the element concerned is given by:

CN
mHQ ——(l°§ —dE)xF‘ xF

 

isc
—°—"g(I°%dE)xF;ij

E 0

I
= C=Co(—)X(ZAF)
I

0

a

( glamxp; xF’

 

where(ZAF)=
dE)xF‘ xF’

hm “Hem gm chow
talks

F ' and F f are factors of absorption and ﬂuorescence for specimen.

F g and F of are factors of absorption and ﬂuorescence for Standard sample.

(ZAF) can be separated into three correction factors which represent effects by atomic

number, absorption and secondary ﬂuorescence.

Z (atomic number) Factor Correction The Z factor term is a combination of S

(stopping power) and R (backscattering) . From Bethe's law S is represented by [46]:

s =-7.85x104(Z/A)(1/E)In(1.166E/J)

44
where J is the mean ionization potential of element concerned. Thus, S is a function of
electron energy E and atomic number. R represents the fraction of the incident electrons'
energy remaining in the sample without loss by backscattering that can produce ionization.
From RuSS'S proposal [47]:
R = e*"°‘Z(U -1)°'3

where U is the overvoltage ratio (deﬁned as E/Ec, E, is critical excitation energy)

Thus, R is a function of Z and U. With both S and R calculable, the total Z term
becomes the integral of (Ra/So) for Standard divided by the integral of R/S for the unknown

matrix.

A (absorption) Factor Correction Whenever X-rays pass through anything,
they are absorbed according to Beer's Law [43]

_ “W
If —I‘.e

where I I is the intensity left, I i is the initial intensity, p is the density and x is the

distance, and p is the 'mass absorption coefﬁcient'.

AS shown in Figure 2— 19 [43], the distance traveled in the specimen by X—rays
generated at a depth 2 is z xcosec( 0), where 0 is the X-ray take off angle. Therefore, the

factor by Which the X-ray intensity is reduced is e-" an “op:

. Hence, deﬁning 4)( pz) as a
function of the depth distribution of X-ray generation, the total absorption factor is an

integral:
Foo = (I)4>(pz)e"“”d(pz)

where X = ttcosece
F‘ = F(x)
Then the total A term correction is just F( x ) for the standard divided by F ( x ) for

the element concerned in the matrix.

45

To Dectector

0 = Take-off angle

 

Absorption
Distance
2 / sin (0)

 

    
 

A\\\\\\'¢\\\\\\

Figure 2-19 Schematic Diagram of X-Ray Generation in Sample and the Absorption
of Photons from Each Depth 2 Enrout to the Detector, Depending on the X—Ray

Take-Off Angle 0 [43]

46

F (fluorescence) Factor Correction In binary systems, element A can be excited
by characteristic ﬂuorescence emitted by element B in the matrix, which will add to the
intensity from element A produced by direct electron beam excitation. This process is
called secondary ﬂuorescence and a ﬂuorescence correction is needed.
A correction for the additional ﬂuorescence intensity from element "1'" due to
radiation from element "1'" has been developed by Reed [48]:
I i

—‘— = 0.5Cjt0j(u‘. lum )((r‘. —1)/r‘.)(A‘. IA). )Pj‘.((Uj —1)/(U‘. -l))5’3
1.

where I f is the ﬂuorescence intensity which is added to the primary intensity I _ , to]. is the
ﬂuorescence yield for element j, r‘. is the absorption jump ratio for element 1' (which is

defined as the mass absorption coefﬁcient on the high energy side divided by that on low
energy side), u, is the mass absorption coefﬁcient, um is the matrix mass absorption
coefﬁcient which is a sum of the coefﬁcient for each element in the matrix times the
element's concentration, A; and Aj are the atomic weights of element 1' and j, the P factor
which is 1.0 for K—K and L-L ﬂuorescence but is taken as 0.24 for K-L and 4.2 for L-K

respectively, U is the overvoltage.
1'
Therefore, the ﬂuorescence correction is (1 + 2 . 4 ), or, 1 plus the sum of the
1*! I ,
fractional additional intensities produced by each individual element in the matrix. Note
that there is no ﬂuorescence in the pure Standard, so no ratio to the pure standard is needed.
Hence if we can calculate Z, A and F, the true concentration of the specimen can be
determined.
But since every correction term depends in one way or another on the composition
of the matrix, such as determining the sum terms of total stopping power, secondary

ﬂuorescence or mass absorption coefﬁcients of the overall elements in the matrix, the

concentrations in the samples must be approximately known in advance. This circular

47
process needs an iterative calculation, carried out by making a ﬁrst approximation to
concentration, computing Z, A and F, applying them to get better concentration values, re-
computing the Z, A and F terms using this second approximation, and so on, until, a

desired levels of accuracy is attained.

CHAPTER 3 EXPERIMENTAL METHODS

The principal goals of this work are to study the effect of the processing
parameters, such as working pressure, voltage, power, and distance, on the composition in
TiNi thin ﬁlms prepared by the dc magnetron sputtering method. These experiments were
undertaken to develop a better understanding of the factors affecting the composition shift
between the target and the deposited ﬁlms. This chapter describes the detailed procedures
used to fabricate TiNi thin ﬁlms and the characterization methods employed to study their

composition and thickness.

3.1 Materials
- The sputtering target alloy, with a composition of Ni51,3Ti43_7, was purchased

from Special Metals Corporation in the form of a UDIME’I'® Nitinol bar, hot rolled and
peeled, with As = -14 °C, Af = -10 °C, 305 mm long and 76.2 mm in diameter. This binary
alloy bar was then machined into round disks, 50.8 mm in diameter and 3 mm thick by
EDM. The surfaces of these disks were then ground by 0.3 um abrasive A1203 particles
and used as targets for the TORUS®-2C dc magnetically enhanced sputtering cathode
assembly.

1000-series aluminum alloy sheets were widely employed for parts in the ﬁxture
design, such as sputter shutter, sputter shielding, and substrate holder. After machining,
these aluminum parts were ﬁrst cleaned in an etchant, consisting of one part HF, three parts
of HNO3 and ﬁve parts H20 by volume, followed by distilled water, acetone and
methanol.

TiNi thin ﬁlms were deposited upon microscope slides (76.2 mm x 25.4 mm). The
slides were used as substrates in the sputtering processes because they facilitated the

process of acquiring free-standing thin ﬁlms for composition analysis.

48

49
3.2 Deposition Equipment

3.2.1 System Geometry

The sputtering processes were carried out in a vacuum system equipped with a
single sputtering cathode, shown schematically in Figure 3-1. The source was mounted
through a standard 25.4 mm hole on the side of vacuum chamber in a 152.4 mm knife-edge
ﬂange. Being in a vertical orientation, neither the substrate nor target can be contaminated
by falling particles. The stem of the sputtering source was modiﬁed by mounting a
threaded aluminum tube (254 mm long and 25.4 mm in diameter), and equipped with a
port-sealing mechanism to prevent leaking on the original short threaded stem. By means
of this modiﬁcation, the processing distance between target and substrate could be directly
adjusted from outside of the chamber without venting.

The sputter shielding with a window (102 mm x 51 mm) opened at the sputtering
position, was ﬁxed between sputter shutter and substrate holder. The shutter plate, whose
vertical position was changed via the moving stage controlled by a rotary feedthrough on
the backside of the chamber, was set at 25.4 mm in front of the substrate holder. The
substrate holder allowed a maximum of ﬁve separately exposable substrates to be installed
at one time. During deposition, the object substrate was moved to the position facing
directly in front of the target by the other moving stage, and the distance between them was
controlled by the movable stem on the sputtering source. Substrates other than the one to
be processed were covered and protected from the sputtered ﬂux by the sputter shielding.

The thickness monitor crystal (See Section 3.2.4) was set at a position directly
behind the substrate to be processed. Intermittent deposition-rate measurements were made

by temporarily lowering the object substrate to expose the oscillator crystal.

50

 

[Bell Jar Vacuum
S

/ _ yStem

 

 

 

 

 

 

 

 

 

   
 
  
 

 

 

    
   

 

 

 

 

 

— ‘ r/Nixri Tar et Movable Stenﬂ
Substrate
Movin Sta e H019.“
Thickness <—>

 

 

 

 

i Monitor ‘
. Substrate ,

D.C. Magnetron Sputtering

Source Assembly

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

\
Shielding

 

 

 

 

 

 

 

 

Eiffusion Pump]

 

Figure 3-1 Schematic Diagram of DC. Magnetron Sputtering System

5 1
3.2.2 Pumping System and Gas Pressure Monitors

The high-vacuum pumping system included one diffusion pump and one
mechanical pump connected in series. The diffusion pump used for the vacuum system is a
model HS-lO made from the Varian/Vacuum Company with a pumping speed of 4000
liters/second of air. A water-cooled trap was used. With the help of low-vapor-pressure
silicone ﬂuid (DC 705), the base-pressrue was reach 8.5x104 Pa without baking the
chamber. After baking for over 12 hours (using the ﬂexible electric heating tapes wrapped
around the chamber wall) and pumping down at least 24 hours, the ultimate base pressme
reached 3x105 Pa with typical partial pressures: P1120 = 1x105 Pa, PN2 = 9x10‘6 Pa, and
P02 = 2x10-6 Pa.

A Pirani Vacuum Gauge from Kurt J. Lesker Company was installed to read the
higher range pressures from 0.1 to 10 Pa and to monitor the pressure dtuing sputtering
processes. An ion gauge was also installed to obtain the lower pressure values such as
base pressure. Furthermore, a Dataun quadrupole mass spectrometer from Spectramass
Inc. was attached to the chamber to measure the partial pressure of residual gases. It also

served as a high-vacuum gauge and a leak detector.

3.2.3 Sputtering Cathode assembly and Power Supply

The construction of the TORUSQ-ZC magnetron sputtering source is illustrated in
Figure 32 Individual segments of the magnet are mounted between circular metallic pole-
pieces and the whole is formed into one solid unit. The cooling water passes up through
the magnet assembly, around its top surface immediately behind the solid copper mounting
jacket where the target was secured by the retaining ring, and down again. Water also
passes down around the outside of magnet assembly, then behind its base, and out the

back.

52

m W-[I—

 

 

 

 

 

 

 

 

 

 

 

 

 

 

] l

a

 

 

 

DARK SPACE SHIELD

THUMB SCREW

HOLD DOWN SCREW

TARGET HOLD DOWN
TARGET

COOLING PLATE SCREW

COOLING PLATE

'O' RING

MAGNET ASSEMBLY

BASE PLATE

POWER TERM SCREW

'0' RING

TEFDON ISOLATION

'0' RING

CATHODE BODY
MOUNTING FLANGE

TEFLON SLEEVE
FLAT WASHER

MTG FLANGE SCREW

'0' RING

WASHER
HEX NUT

Figure 3-2 Exploded View of TORUS-2C Magnetron Sputtering Source.

53

The TORUS®-2C sputtering source design is based upon a modiﬁed Penning
Discharge Principle. This design, with its powerful Samarium-cobalt magnets and shaped
pole-pieces enhances the magnetic ﬁeld immediame above the plane of the target. Instead
of allowing secondary electrons that are emitted during sputtering processes to impinge
upon the substrate, TORUSO-2C can capture electrons released near the target, concentrate
them, and employ them to develop higher sputtering efﬁciency.

The power supply was a model Advanced Energy MDXIK. The unit will deliver 1
kW to a magnetron cathode at 1000 V and is conu'olled with any of three regulation modes:

power, current or voltage.

3.2.4 Thickness Monitor

A quartz crystal thickness monitor (Leybold Inﬁcon Inc. XTM/‘Z Deposition
Monitor), consist of a quartz crystal sensor (Standard Sensor 750-207-Gl) and a
monitor/controller unit, was employed to measure deposition rate. The quartz crystal was
fabricated into a disk with 12 mm in diameter, and had a 6.0 MHz resonant frequency. It
was held right behind the substrate holder at the processing substrate position. Since
measurements were intermittent, the quartz crystal sensor's water cooling system was not
used.

In principle, the resonant frequency, f q, of the quartz crystal sensor varies as a
function of the mass of matter deposited on it [49]. As a result, the mass of ﬁlm deposited

can be expressed as
at = (N,p,, Inf,2){tan“ [Ztan(n(fq — f, w, m

where M, is the frequency constant for AT-cut quartz crystal, p q is the density of quartz,

f q is the resonant frequency of uncoated quartz, f c is the resonant frequency of coated

quartz, and Z is the acoustic impedance ratio which is written as

54

)1/2

Z =

(13qu I Pill,

where p f is the density of quartz crystal and deposited ﬁlm, and ttq and ttf are shear
moduli of quartz crystal and deposited ﬁlm respectively. The deposited ﬁlm thickness can
be calculated ﬁom ﬁlm mass and ﬁlm density accordingly. In the present work, the density
of alloy ﬁlm is replaced by its bulk density of 6.5 mm, and the shear modulus replaced
by 2.48x10" dyne/cmz. The p q and ttq- are obtained from manual 2.649 gm/cm3 and

3.32><1011 dyne/cmz. So, the 2 factor is calculated to be 0.7386.

3.3 Deposition Procedure

3.3.1 Preparation

Substrate Cleaning All substrates were cleaned by acetone and methanol,
respectively, in an ultrasonic cleaner for about 15 minutes to remove grease contamination
prior to being installed. the substrate holder was polished by sand paper (grid No. 600) to
remove previously deposited coatings, and then cleaned by distilled water, and
subsequently by acetone and methanol.

Evacuation The vacuum chamber was evacuated by diffusion pumping
for more than 24 hours after the pressure was brought down to 4 Pa by the mechanical
pump. Electric heating tapes were used to bake out the adsorbed water vapor yielding the
vacuum system base pressure below 3x105 Pa

Pressure Set Up The chamber was back-ﬁlled with argon gas and the Mass
Flow Meter/Mass Flow Controller (from MKS Instruments Inc.) was used to control the
chamber pressures in the region from 0.4 to 1.1 Pa. Detailed process conditions will be
listed later.

Target Wear-In A N in 3Ti43,7 alloy target was used to deposit the ﬁlms.

During installation of the target, exposure to some physical handling and to atmospheric

55
gases was unavoidable. Where the goal is to deposit ﬁlms composed exclusively of the

pure target material, target wear-in/clean up should be done prior to deposition on the
substrate behind a sputter shutter. The virgin alloy target was pre-sputtered for 0.5 hour at
higher power levels to attain steady-state sputtering conditions. This procedure removed
oxides and contaminants and prepared it to produce high quality ﬁlms. This preliminary
wear-in also cause thermal expansion of the target which seats it properly in its cooling
mount in preparation for later runs. Afterwards, the alloy target was pre-sputtered for 10
minutes under the experimental process condition at the beginning of each individual

deposition run.

3.3.2 Processing

Condition Set Up There were ﬁve process factors considered in this
study: target erosion, working pressure, cathode voltage, cathode power and working
distance. For experimental convenience, only one variable was held constant per test. The
detailed process conditions for every test run will be shown in section 4.2.

Deposition Rate Monitor At the beginning of each experiment run, the whole
substrate holder was at a position lower than the processing position. After the target was
pie-sputtered for at least 10 minutes under the desired process condition, the thickness
monitor was exposed by raising the sputter shutter. Thus, the thickness monitor began to
measure the deposition rate. Exposure of the sensor to the sputtering source was limited to
one minute to avoid a temperature increase in the quartz crystal, however, the exact
temperature of the quartz crystal was not known. The deposition rate was assumed to be
constant during a run.

Film Deposition Prior to exposing the substrates to the sputtering
source, the deposition rate was estimated by the thickness monitor, and the desired

sputtering time was determined such that the total thickness of ﬁlm was at least 3 pm

allowing for the accurate measurement of composition. A series of substrates was set on

56
the holder. One substrate was exposed per run. When the processing conditions
stabilized, the target shutter was opened allowing for the deposition of thin ﬁlm.

3.4 Composition Analysis

Elemental composition of the as-deposited ﬁlms was determined by energy
dispersive X-ray spectroscopy (EDS) microanalysis in a Hitachi S-2500 scanning electron
microscope (SEM). This technique can determine the presence of various elements by
analyzing the different energies of their characteristic X-rays, which are indirectly
stimulated by incident electrons. Although the probe of incident electrons is very small, the
area where X-rays are produced is much wider than the electron probe and reaches deeply
into the sample. Films with thickness of at least 3 um on glass substrates were illuminated
by the electron beam with an energy of 20 keV. Since no Si characteristic radiation was
observed, the substrate was taken to have no interaction with electrons, and the spectra
could be then treated by a ZAF correction.

A program called ZAF~4, supplied by Link Analytical Inc., was used to analyze all
the spectra, which were collected under carefully controlled conditions. The optics of the
electron beam and the geometry between lens, sample and X-ray detector were rigorously
kept consistent during data collection. The LaB6 electron source has to be stabilized for at
least 0.5 hour prior to taking any spectra. All the spectra were collected at the same
working distance (15 mm), magniﬁcation (2000 x), tilting angle, dead time percentage and
time interval.

The electrical parameters of the lens remained untouched throughout the whole
process and the height of sample was adjusted to maintain a constant working distance.
The Ka peak, acquired from a pure nickel standard and employed as a calibration
spectrmn, was collected every 15 to 20 minutes. Pure titanium and a Ni49,9Ti50_1 alloy
were employed to be the standards for composition analysis, and the composition of target

alloy was also measured simultaneously to evaluate the accuracy of this method. The bulk

57
materials were last polished by using 0.05 pm alumina powder to yield mirror-like
surfaces, and then cleaned by acetone and methanol. The specimens, including standards,
were loaded on a graphite stage without tilting.
The composition of the sputter-target alloy determined by the ZAF method was:
50.59 at% Ni-and 49.41 at% Ti. The results shows that the measured titanium content is
about 0.7 at% higher than the manufacturer's speciﬁcation, which was considered to be

nominal.

3.5 Thickness Measurement

The thickness of the as-deposited ﬁlms were also used to verify the measured
deposition rate. To obtain a thickness proﬁle, a Dektak II proﬁlometer was employed to
scan across a 3 mm wide step produced by masking the substrates during deposition. The
proﬁling instrument draws a ﬁne stylus along the surface containing the stepped ﬁlm with
speed 1.5 mm/min. Whenever the stylus encounters a step, a signal variation (based on
differential capacitance) yields an indication of the step high. The information is displayed
on a CRT.

CHAPTER 4 RESULTS AND DISCUSSION

Druin g the deposition of alloy ﬁlms, the various factors that would cause the
composition of sputtered ﬁlm to be different ﬁom that of the target alloy composition were
generally described in section 2.4.5. In this study, TiNi thin ﬁlms were fabricated by dc
magnetron sputter-deposition under different process conditions. Actually, processing
parameters were chosen in order to determine which factors mentioned above would
dominate during the deposition process. Therefore, there are two major parts in this
chapter. The ﬁrst part concentrates on the discussion of how these factors can be applied
to this study. Then, in the second part, the experimental results showing the effects of
sputtering process parameters on the composition of sputtered ﬁlms are presented and

interpreted by means of the discussions reported in the ﬁrst part

4.1 Factors Determining The Composition of The Sputter-Deposited Thin Films

For metallic targets, the sputtering processes can be described by elastic collision
in the surface layers of a solid. If an incident ion with sufﬁcient energy collides with an
atom of the solid, it creates a primary knockon atom. For the collision cascades started by
this primary knockon atom, it is convenient to distinguish between three qualitatively
different situations [50].
(l) The single-knockon regime. Recoil atoms ﬁom ion-target collisions receive
sufﬁciently high energy to be sputtered, but not enough to generate recoil cascades.
(2) The linear cascade regime. Recoil atoms ﬁom ion-target collisions receive sufﬁciently
high energy to generate recoil cascades. The density of recoil atoms is sufﬁciently low so
that knock on collisions dominate and collisions between moving atoms are infrequent.
(3) The spike regime. The density of recoil atoms is so high that the majority of atoms

within a certain volume are in motion.

58

59
Qualitatively, the sin gle-knockon regime applies for ion bombardment at low

energies, i.e. close to the sputtering threshold and for light ions, where mostly only small
energies are transferred to target atoms. The linear cascade regime characterizes ions with
medium to large atomic number in the keV energy region. The spike regime is reached for
bombardment of large atomic mass targets with large mass ions or molecular ions at
energies around 20 to 80 keV.

Since the planar magnetron source is usually operated in argon at a pressure of
0.15 to 1.5 Pa and at cathode potentials of 300-1000 V, under these conditions, the model
for sputtering is limited to the linear cascade regime. However, there is no clear boundary
between linear cascade regime and single-knockon regime, and the single-knockon model
will be invoked as necessary.

4.1.1 Preferential Sputtering And Component Sputtering Yield Ratio

Due to preferential sputtering in the sputtering of alloy materials, the preferential
removal of one component from the surface leads to the formation of the so-called "altered
layer", which is a near-Stu'face region with composition different from the bulk
composition. At sufﬁciently low temperatures, when thermal diffusion is not important,
the altered layer stays at a ﬁnite thickness and steady-state conditions can be reached
where the amount of material sputtered from each species becomes proportional to the
bulk concentration.

At steady-state, the ratio of the number of sputtered atoms n, of component A to n,
of component B must be equal to the ratio of the bulk concentrations cA / c ' [41]

nA/nn=cA/c, (4.1)
Since the ratio of the number of sputtered atoms n A / n5 is equal to the ratio of the

product of component sputtering yield Y,- and equilibrium surface concentration c: of

component i during sputtering, then

60
s s
0.4/cs =nA/nn =YA 'cA/Ya '63

= 1‘— = 573/53 (4.2)
YB CA /C'

Equation (4.2) allows the experimental determination of the ratio of the component
sputtering yields by measuring the surface and as well as the bulk composition, i.e. the
changes in surface composition under ion bombardment. In addition, the component
sputtering yields can be compared to the yields of the ptne elements. Generally, the
sputtering yields for the alloy target components are different from those of the pure
elements.

Based on Equation (4.2), systematic studies on alloy composition were performed
on a number of binary alloy systems [41]. The experimental results showed that the
observed surface enrichment is mostly in agreement with the sputtering yields of the pure
elements, i.e. the component with the lower elemental yield becomes enriched under ion
bombardment.

Different models have been proposed to predict the changes in surface
composition, such as mass effects and surface binding effects [50, 51, 52]. The mass
effects model predicts that the depletion of the lighter target component and it becomes
prevalent with decreasing energy. The surface binding energy argument predicts the
enrichment of the component with the higher surface binding energy or heat of
atomization. The different effects can give opposite results and generally a combination of
these two effects will occur. However, it was found that enrichment after the treatment of
mass effects and surface binding effects is in most cases too small as compared to the
experimentally observed changes [52].

In the case of TiNi alloy system, properties comparison between the nickel and
titanium of atomic mass, atomic radius, heat of atomization, sputtering threshold energy,

and the sputtering yields at 600 eV is indicated in Table 4-1.

 

61
Table 4-1 Properties Comparison of Titanium and Nickel

 

 

 

 

 

 

 

 

 

Titanium Nickel
Atomic Mass 47 .90 58.71
Atomic Radius 1.46 A g 1.25 A
Heat of Atomization 473 KJ/mole 429 KJ/mole
Sputtering Threshold Energy 20 eV 21 eV
Sputtering Yield (at 600 eV) 0.41 atoms/ion 1.5 atoms/ion

 

Since pure elemental Ni has a higher, sputtering yield than pure elemental Ti under
the same ion bombardment energy, it is believed. that YN, > Y“ is still true in the
sputtering of TiN i alloys. From the mass effects model [50], the component with lighter
mass and lower threshold energy is in favor of being preferentially sputtered and becomes
more prevalent with decreasing energy. This means that the sputtering yield ratio YNil Yvr,
will decrease as the energy decreases. _

The component sputtering yield ratios in the magnetron sputter-deposited TiN i
amorphous thin ﬁlms were studied by Chang using IBAD method [26]. The experimental
results showed YT, > YNi and the ratio of Y-n/ Ym (note inverse ratio) increased from 1.75
at 500 eV to 9 at 50 eV. The trend for the changes of component sputtering yield ratios is
the same as that for crystallized alloys, and can be explained using the same mechanism.
However, Y“ > YNi in the amorphous alloy is opposite to the assumption of YN; > Y—n in
the crystalline alloy. This phenomenon could be understood considering the surface
binding energy effects. For a crystalline binary alloy AB, the surface binding energies UA
and U; have been expressed by Kelly [52] as functions of the nearest-neighbor bond

strengths UM, U33 and U“. Thus, one obtains for UA and U3:

S

S
UA = —CAZSUM — CBZSUM (4.3)

62

U -c 2 U -c‘z,U (4.4)

3535 AB

where Z, is the surface coordination number, and c: is the equilibrium surface

concentration of component i drning sputtering. As for amorphous alloys, the structures
almost have a complete lack of periodicity and there are no speciﬁc distinctions between
UM, U33 and U“. All bond strengths can be assumed identical in the amorphous alloy.
Therefore, there are no differences between UA and U3, and the mass effects dominate the
preferential sputtering. This will cause the depletion of Ti component in the amorphous
TiNi alloy during sputtering. Because the as-deposited TiNi thin ﬁlms are in amorphous
form, Ni—enrichment thin ﬁlms are predicted due to the preferential resputtering of Ti

component.

4.1.2 Angular Distribution of Sputtered. Atoms

In the case of sputtering by linear cascades the ejection distribution from a
polycrystalline solid is largely independent of mass, energy and incident angle of the
bombardment ions [53, 54]. The ejection distribution is rotationally symmetric with

respect to the surface normal and conforms to [54]

dY Ve
-— = cos 0
d!)

e 1 5 ve < 2 (4.5)
where Y is sputtering yield, {2 is the direction of emission, and 0., is the polar angle of the
emitted particle. A more or less pronounced "over cosine" distribution (vc > 1) is often
observed [55]. Such behavior is expected from the mechanism of linear cascades and is
reproduced by computer simulation [56].

In the low-energy regime of medium to heavy ion sputtering, i.e. m > 4 and E < 1
keV, the emission distribution depends on the incident angle, and deviates from the cosine

law giving higher emission at larger polar angles. These so called "under-cosine"

 

63

distributions are generally explained as a result of shallow cascades, where too few
collisions are available to acquire momentum randomization [53].

Then, if an alloy is sputtered, the angular distribution of the sputtered atoms can be
different for each component. Basically, they are all assumed to follow the cosine law
distribution but have different shapes, as shown in Figure 41 [57].

Over Cosine Surface Normal

Under Cosine

Cosine

 

 

 

Figure 4-1 Angular Distribution of Sputtered Atoms from a Polycrystalline Target [57]

64

In a study by Olson [5 8], for normal ion incidence and different ion energies, Fe-
Ni and Ni-Cu alloys were sputtered and the emitted atoms were collected at various
angles. The results showed that at lower energy (< 100 eV), preferred ejection of the
lighter mass atoms in the target normal direction occurred. The phenomenon was not
unexpected since, when only a few atom collisions become involved, the fact that a lighter
mass atom can be backscattered from a heavier one, but not vice versa, must reveal itself
in the observed preferential ejection of lighter masses in the upstream ion beam direction.

However, at higher energies (200-1000 eV), the experimental results were
contradictory. Sputtered Fe-Ni alloy led to a preferred normal ejection of heavier
components, while for Ni—Cu alloy, there was no change in preferred normal ejection of
the lighter component. Thus, for higher ion energies (> 200 eV), the results indicate that
factors other than mass, such as binding energy and crystal structure, become dominant
and determine the angular distribution of the component. This effect implies that in ﬁlm
deposition at low gas pressure, the composition depends upon the geometry and relative
position of the target and substrate.

The geometric relationship of substrate and target surface in this study is shown in
Figure 4-2, there are two collection positions on the substrate. One is at the position of
the target center axis, and the other is at a radius of 2.5 cm, angle zero degrees in the polar
coordinate system relative to the axis of the target. Note that the eroded target surface
angle is measured from a target that has been sputtered for about 15 hours, as shown in
Figure 4-3. Assume that the distribution of the sputtered atoms from the inclined target
surface follows the cosine law with respect to the surface normal. Then, referring to
Figure 4-2, it shows that the original preferred normal ejection constituent will be
preferably collected at the center position of the substrate, and this effect will increase as

the target erodes.

65

 

Center Off Center
POSIUOII Position
Substrat
17° 1 1°
25°
Eroded Target

Figure 4-2 The Geometric Relationship of Samples and Target Surface

 

Figure 4—3 An Eroded Target with its Localized Erosion Angle

66
4.1.3 The Development of Target Surface Topography

A topography will deve10p upon heavily ion irradiated surface [59]. Sputtering
creates at least atomic scale discontinuities (< 100 A) at the surface, and such effects may
be observed for individual ions and for ﬂuences up to the order of 101‘ ions/cmz,
depending on ion species and target materials. When ion ﬂuence increases above the
order of 1017 ions/cm2, the density of atomic scale discontinuities increases, and the
bombardment induced defect features become microscopically observable and range in
sizes of 100-10,000 A. Then, local variations in sputtering yield occtu resulting in major
changes in surface topography, as evidenced by the development of topographic features
such as etch pits, cones and ripples. At a very large ion ﬂuence, i.e., above the order of
1019-1020 ions/cmz, sputter etch effects may induce differential dimensional changes in the
order of sub-millimetric size. At such levels of dimensional change, theoretical models of
morphology change become meaningful.

In this study, the sputtering target had a bombardment area of about 41: c1112 and
the ion current ranged from 60 mA to 500 mA. Therefore, the ion current densities ranged
from 5 to 40 mA/cm2 and in other units from 3.15x1016 ions/cmzsec to 2.5x10”
ions/cmzsec. Since the deposition time was at least 20 minutes for every test run, the
ﬂuence was above the order of 10‘9 ions/cm2 for every process condition. Thus, the
macroscopic theory is applicable to this study. Figure 44 indicates the grains and grain
boundaries of a target before ion bombardment. Figure 45 and Figure 4-6 show the
topography of the target surface after about 15 hours ion bombardment. Ridges and cones
can be easily found in the ﬁgures.

67

 

 

 

Figure 4-4 The Grain Structures of Target Surface before Ion Bombardment
(100 x)

68

 

 

 

Figure 4-5 The Topography of a Target Surface Sputtered for 15 Hours

69

 

 

 

Figure 4-6 The Cone Structure on the Target Surface

70

A ﬁrst-order erosion theory, ignoring secondary effects such as local ﬂux
variations, atomic redeposition and surface atomic transport, was proposed by Carter and
Barber [60, 61], and successfully predicted the formation of topographic features, in
particular cones, which have been experimentally observed on the macroscopic sputtering
scale and at the microsc0pic level. The basic assumption of this theory is that sputtering
yield varies with the incident angle and the angle of maximum sputtering yield is present
within initial protuberances above the mean surface. For initial low roughness surface,
the high angle surfaces required for this theory can form ﬁom either contaminant
protection, as shown in Figure 4-7 [62], or from grain boundaries, etch pits and any other
departme from ﬂatness. If the initial surface is itself rather rough, then ridge and conical
structures may develop from the beginning.

Also, the development of surface topography is sensitive to the bombarding ion
current density. For example, the surface of a copper crystal can become covered with
cones when the ion current density is > 100 itA/cmz, but, with the same total ﬂuence with
a ion current density of < 100 ttA/cmz, only occasional cones are produced [59].
Fmthermore, Kaufman and Rossnagel [63, 64] reported the average cone spacing is
proportional to the inverse square root of the bombarding ion current density. It means
that a higher ion current density causes a higher density of cones on the bombarding
surface.

it

71

 

Contamination or Impurity

 

 

 

 

 

 

 

 

 

Figure 4-7 Schematic Diagram Showing Impurity-Induced Cone and Subsequent Pit

Formation [62]

 

72

An important phenomenon related to surface topography is the variation in total
sputtering yield. Because of the increase in yield caused by oblique incidence (Figure 2-
12), the feature structured surfaces lead to enhanced emission; but recapture of ejected
particles at these ridges or cones may compensate for this effect. Littmark and Hofer [65]
considered this problem and found the total sputtering yield to be larger than that for ﬂat
surfaces. Only in the case of steep, needle-like cone structures (incidence angle > 70°)
will the recapture of released particles prevail and the yield be smaller. However, there are
no studies concerning the relationship between surface topography and the changes of

component sputtering yield in the sputtering of alloy .

4.1.4 Resputtering of Growing Thin Films

Figure 2-15 shows that the substrate is inevitably subjected to impingement by
many species during the deposition process, such as ions and neutral atoms, so there
always exists a certain amount of resputtering of the growing ﬁlm.

The source of neuual atoms is the primary bombarding ions reﬂected from the
target surface. Vossen [66] reported that these bombarding ions have a high possibility of
being neutralized prior to impacting the target. Then, they may be reﬂected as neutral
atoms, not ions, without being inﬂuenced by the electric ﬁeld over the target surface. The
amount of reﬂection is an inverse function of primary bombarding energy because this
effect competes with ion implantation. At low primary energies (< 1000 eV), reﬂection
fractions as high as 0.4 have been observed, whereas at high primary bombarding
energies (> 1000 eV) typical reﬂection fractions are of the order of 0.05 [66]. However,
the reﬂected atoms have a wide energy spectrum which make it difﬁcult to estimate the
incident ﬂux of these particles.

In magnetron sputtering, the magnetic ﬁeld does not directly affect the ion motion.
Nevertheless, because of electrostatic attraction, the ions move with the electrons keeping

the plasma neutral. Theoretically, in a conventional magnetron, most of the discharge is

73
conﬁned close to the cathode surface and, therefore, bombardment of growing ﬁlm by

electrons and ions is minimized. In practice, substrate bombardment could be
signiﬁcantly increased by "unbalancing" the magnetics [67]. In an unbalanced magnetron.
the ﬂux from the north pole is unequal to that entering the south pole, as indicated in
Figure 4-8 [68]. The ﬂux from the central is less than that of the outer magnet.

Therefore, the system gives large ion currents and large electron currents to the substrate.
Electrons are channeled along ﬁeld lines extending from the discharge region to the
substrate. The unbalanced magnetron is capable of giving ion ﬂuxes at the substrate that
are larger than the ﬂux of sputter atoms. The resputtering ion ﬂux is very dependent on
the magnetic ﬁeld conﬁguration, discharge current; however, it is roughly independent of
the target composition, gas pressure, and gas composition [67, 68].

Because magnetic ﬁeld lines are difﬁcult to focus, it is very difﬁcult to construct a
perfect, "balanced" magnetron. Thus, magnetron sources are always accompanied by
substrate bombardment by ions and electrons from the plasma discharge.

According to the IBAD results by Chang [26], the sputtering yield of Ti in the
deposited amorphous TiN i thin ﬁlm is higher than that of Ni (YT, > Ym). Therefore, Ti
will be preferentially removed from the ﬁlm. Furthermore, unlike the target, the sm'face
of the growing ﬁlm is constantly being replenished with new materials of the cathode
composition, so an "altered layer" cannot form and the "steady-state" cannot never be
established. Thus, the depletion amount of Ti will increase as the resputtering effect

increases.

74

 

 

 

 

MAGNETICS

 

 

 

 

 

 

 

Figure 4-8 Schematic Representation of an Unbalanced Magnetron [68]

CATHODE POTENTIAL (V)

75

 

 

       
  

 

 

AVERAGE CATHODE CURRENT DENSITY (mA/cmz)

Figure 4-9 A Typical Voltage-Current Characteristic for Planar Magnetron
Cathode at Various Pressures [69]

700 r r T , T I I j
600 - ARGO” 1m'rorr 2mTorr -
500- l
400” 4
300~
IOmTorr

200- .
’100 a . a . . . I 4 . . ' . . .

4 6 8 10 20 4o 60 80

76

4.1.5 Planar Magnetron Source

The typical voltage-current characteristics for a planar magnetron source is shown
in Figme 4-9 for various pressures. These curves have been empirically found to follow
the relation [69]

1 = kV " (4.6)

where I is the cathode current (or current density) and V is the cathode voltage, k is a
constant, and the exponent n is a constant ranging ﬁom 3 to 15. The constants of k and
the exponent n are dependent on the type of magnetron used, the target materials, the gas
species, and the gas pressure. Usually, the more efﬁcient the electron trapping in the

plasma, the higher the exponent n.. Also, since the power (W) is the product of cathode

voltage and current, Equation (4.6) becomes

W = [V = kV"+1 (4.7)
where k and n are the same constants as in Equation (4.6).

Thus, for a given target material and cathode conﬁguration, a family of current-
voltage curves versus pressure can be obtained. This enables a suitable operating point to
be chosen for a given voltage or power value.

However, as the target erodes, the plasma impedance changes because of the
increased strength of the parallel component of the magnetic ﬁeld at the surface of the
target [70]. This indicates that the exponent n in Equation (4.6) will increase as the
process continues. For most applications, this does not pose a problem. But, in the
present experiment, the magnetron source used for sputtering is relatively small (5 cm in
diameter), and the change of the exponent n causes an evident change of the process
condition during the deposition. This may induce other factors that have to be considered
for the control of composition of the deposited thin ﬁlm.

As mentioned in section 3.2.3, in this study the deposition process can be
controlled at power, voltage, or current regulation mode, meaning that power, voltage, or

current is kept at a constant throughout the deposition periods. In magnetron sputtering,

77
from Equation (4.7), the k and exponent n both are constants at a speciﬁc process
pressure, and controlling the process by power or voltage has the same effects. But,
actually, the exponent n will increase during deposition even though under the same
pressure. Thus, when the process is controlled by power regulation mode, the voltage
will decrease with increasing exponent n to maintain a constant power; when the process
is controlled by voltage regulation mode, the discharge current increases as the exponent n

increases.

4.2 Effects of Process Parameters on Thin Film Composition

There are four process parameters being surveyed in the present experiment, they
being plasma voltage, plasma power, process pressure, and working distance between
target and substrate, respectively. Before these tests for the process parameters, a test was
conducted to determine the signiﬁcance of target erosion effect on the thin ﬁlm
composition. The ﬁnal composition of the deposited ﬁlm for every process condition is
determined by the combined effects of various factors considered in the ﬁrst part of this
chapter. In addition to the complexity of the glow discharge sputtering environment and
the formation of target surface topography features by high bombarding ﬂuences, the
continuous changes of strength of the magnetic ﬁeld (i.e., the exponent n in Equation
(4.6)) make the process conditions non-steady-state and extremely complicated. Most of
the experimental results can be only interpreted qualitatively at this stage. However,
valuable information concerning the fabrication of TiN i thin ﬁlms using a magnetron
source TORUS®-2C can be obtained.

The sputtering target has a ﬁnite life and will eventually erode through. Also, the
development of topographies on the target surface with processing will affect the
deposited ﬁlm composition. Usually, the amount of target erosion is estimated by the
value of (time x current). Since the deposition processes are normally operated under

1000 eV in magnetron sputtering, the amount of erosion can also be estimated by (time x

78
power) due to the linear increase of sputtering yield with voltage in this energy range. A
table of target histories for the test factors affecting deposited ﬁlm composition is shown
in Table 42. The values recorded in this table are the products of time and power. There
was a total of two targets being used in this study. Approximately, the total target life is

about 3800 W-hour, and it can be used up to 2800 W-hour without eroding through in
the deposition of TiN i thin ﬁlm by magnetron source TORUSO-ZC.

Table 4-2 Histories of Targets in the Different Factor Tests (W hour)

 

 

 

 

 

 

 

 

 

 

 

 

Test Target Used Before Used in This * How Much
Number This Test ‘ Test Left
Target Erosion 1 1000 1050 1750
Plasma Voltage 1 2050 550 1200
Plasma Power 2 0 500 3300
Gas Pressure 2 500 900 2400
Working Distance 2 1400 450 1950

 

* Total target life is about 3800 W-hour and about 2800 W-hour available .

79
4.2.1 Effect of Target Erosion

One disadvantage of using a planar magnetron source is the localized erosion of
the target. This series of experiments is thus designed to study if non-uniform erosion
eﬁects can have inﬂuences on ﬁlm composition. The experimental procedure for this test
was controlled by power regulation mode while keeping the other process parameters such
as pressure and distance constant. Then, thin ﬁlm samples were continuously collected at
different process time periods. The process conditions and their compositional results are
shown in Table 43. Since the processes were controlled by power regulation mode, the
decreasing of voltage and increasing of discharge current during processing were
expected, as predicted by Equation (4.7). However, the quick decrease of voltage during
every deposition was unexpected.

A graph of the target erosion effect on the composition of thin ﬁlms is shown in
Figure 4-10. In this graph, thin ﬁlm with the highest Ni concentration was found at the
beginning of the process, then went to a stable region, and at the fourth hour, a Ti-rich
thin ﬁlm with respect to the target alloy composition was produced.

Assuming that all the process parameters were constant during deposition, the
changes of composition with target erosion time could be attributed to one cause —
changes in target surface topography, macroscopically and microscopically. Due to the
preferential ejection of the high sputtering yield component at the beginning of the
process, the sputtered ﬂux leaving the target surface has a higher concentration of the high
sputtering yield component. The high sputtering component in TiN i alloy is Ni, as
discussed in section 4.1.1. So the ﬁrst thin ﬁlm sample had thehighest Ni concentration.
After the target "wear-in" for approximately 20-30 minutes, a steady-state is reached and
the sputtered ﬂux has the same composition as the target alloy. But, since the resputtering
effect cannot be avoided and causes the depletion of Ti, the deposited thin ﬁlms still
possessed a higher Ni concentration than the target alloy. Finally, the target surface was

built up to an extent which could change the stable region.

80

 

 

 

Cathode Current

Deposition time for each sample = 60 min Distance = 5.08 cm
Power -.- 250 Watts At Center Position Pressure = 0.41 Pa
WWW-a 739v 738 -)689V 687 —>643V I 64w

ﬂ4a336mA340a365mA367—>387mAI388—>400mA

 

 

 

 

 

 

 

 

 

Time Interval 1st hour 2nd hour 3rd hour 4th hour

Thickness 91683 A 87973 A 86088 A 82514 A
—_Dep. Rate 1530 A/min 1465 A/min 1435 A/min 137T min
Composition (M at%) 51.73i0.14% 57)?4:to.22% 51.10:I:O.20% 50———.25i0.12%

 

 

Target Composition (Ni-Ti at%) 50.59 - 49.41 at% :I: 0.04 Ni

 

 

Table 4-3 Process Condition for Target Erosion Test and the Compositional Results

 

 

 

 

 

 

 

 

 

  
 

 

 

 

 

52.0 '1'

51.5 v
A ‘I'
e? .
‘5 \ I __:_,
"z- 51.0 "" ,,____ ..
g I
.a - _ ........ _ - _ ....... - - - ............... Sputter Target
0% 50.5 .0;- ——————— --- ------ 'I - - "‘ --------------- comWSinon
8

50.0 --

49.5 i i i

1st hour 2nd hour 3rd hour 4th hour
Target Erosion Time

Figure 4-10 Composition Change as a Function of the Target Consumption

8 1
There are two effects resulting from the development of surface topography.

Macroscopically, the change of target smt'ace angle will cause a different angular
distribution of sputtered atoms. As mentioned in section 4.1.2 and referred to Figure 42,
this effect may cause the original preferred normal ejection constituent to be preferably
collected. Microscopically, the formation of target surface topography features such as
ridges, cones, and etch pits could change the total sputtering yield and the component
sputtering yield. Because of the absence of experimental data concerning the effect of
surface topography on the component sputtering yield, it is difﬁcult to determine which
component is preferentially ejected. Nevertheless, the experimental results in this study
show that the combination of macroscopically and microscopically topographic effects
produced a Ti-rich thin film.

However, the above discussion is not totally applicable because the process
parameters, such as voltage and discharge current, were not kept at a constant dming
deposition. Thus, in addition to the surface topographic effect, there is another factor that
must be taken into consideration, i.e. if the depositions were always processed under a
steady-state, the sputtered ﬂux leaving the target surface always carried the same
composition as the target alloy.

At the beginning of the process, the sputtered ﬂux had a higher Ni concentration as
did the ﬁrst sample. After about 20-30 minutes, steady-state was established at a speciﬁc
voltage and an altered layer formed with a Ti-enrichment composition related to the
component sputtering yield ratio Ym/Y 1', at this speciﬁc voltage. In a short time (1~2
minutes), the voltage dropped and the original altered layer had to change to a new altered
layer having another Ti-rich composition related to YNi/Y 1", at the second voltage to
establish the steady-state. Since the value of YN-J YT; decreases with decreasing energy,
as discussed in section 4.1.1, the altered layer will become less Ti-rich with decreasing
energy. Thus, the sputtered ﬂux will carry a higher Ti concentration during the transition

period than when it did at steady-state.

 

82
Previous experimental results for other alloys showed that the removal of one to

two altered layer thicknesses was necessary to reach the steady-state condition [41].
Usually, the thickness of an altered layer is determined by the maximum penetration depth
of the primary ions, and a combination of preferential sputtering at the surface and a
transport mechanism. The transport mechanism describing the propagation of the changed
surface composition deeper into the material includes thermal diffusion and/or radiation
enhanced diffusion, and erosion of the target due to sputtering [71]. Radiation enhanced
diffusion means that under ion bombardment, many lattice atoms are displaced in the
course of the collision cascade creating vacancies, interstitials and ﬁnally larger defects, all
of which enhance diffusion. Ho [71] extended a kinetic analysis of preferred sputtering
[72] by including radiation enhanced diffusion. Solutions for the steady—state showed that
the composition of the altered layer varied with the distance from the surface in a simple
exponential manner, and the effective altered layer thickness 5 in the alloy was 5 = D/u,
where D is the diffusion coefﬁcient, u is the velocity of the receding surface (erosion rate).
The expression of 5 was already predicted by Pickering [73]. Also, Ho [71] estimates the
magnitude of diffusivity from measurements [72] yield values of the order of 10“ cmZ/sec
for Cu-Ni alloy. However, since the dc plasrrra sputtering system usually has a high
sputtering rate (>10 A/s), the thickness of altered layer determined by diffusion effect is
not evident.

Ion bombardment of an alloy generally produces an altered layer whose
composition at thermal equilibrium no longer corresponds to a single-phase alloy. Thus,
precipitation of a component cannot be excluded, especially for the TiN i alloys. The
existence of precipitates will cause selective sputtering, based on different erosion rates of
different phases. For high ﬂuences (2 1019 ions/cmz), the interplay of selective sputtering
and the formation of surface topography features ﬁnally leads to a steady-state condition.
For selective sputtering being the dominant process, steady-state conditions will be

reached after sputter erosion of a layer on the order of a few grain diameters, as has been

83
demonstrated for the case of MgO/Au cermets [74]. If topography effects dominate, an

equilibrium is reached after development of a quasi steady-state surface topography, i.e.,
when the rates of cone formation and annihilation are equal. In a study for sputtering Ag-
Co and Ag-Ni, the steady-state conditions were reached after the removal of a layer much
smaller than the average grain diameter [75], indicating the dominance of topography
effects over the selective sputtering. However, no matter which effects dominate,
sputtering of such a solid is marked by a long transition time before steady-state
conditions are reached.

Thus, in this study, when processes were controlled by power regulation mode,
the sputtered ﬂux had high possibility of carrying more Ti than at steady-state due to the
long time transition periods, and the total amount of T1 was decided by the difference of
the component sputtering ratio between energies from the process beginning to the end. If
a function of component sputtering yield ratio Ym/Y-r, versus energy is obtained, the slope
of this function during the energy interval determines how signiﬁcant this effect is.
Therefore, except for the ﬁrst sample, the composition of the other three thin ﬁlm samples
can be explained in this way. Remember that the resputtering effect would cause the
depletion of Ti in the deposited thin ﬁlms even though the sputtered ﬂux is Ti-rich.

' Therefore, there are three competing factors affecting the ﬁnal ﬁlm composition,
the macroscopic and microscopic changes of target surface topograpy, and the non-steady-
state effect. However, at this point, it is difﬁcult to distinguish which one is the dominant

factor in this test.

84
4.2.2 Effect of Plasma Voltagc

For glow discharge sputtering process the plasma voltage (cathode potential) is the
most important process parameter. It determines the energies of the bombarding ions and
is related to the sputtering yields of the sputtered materials. In this test, the process
conditions were controlled by voltage regulation mode and at three different voltage
values: 500V, 625V, and 750V. Both the gas pressure and working distance were kept
constant at 0.53 Pa and 50.8 mm, respectively. Detailed process conditions and their
compositional results are shown in Table 4—4. Notice that the deposition time for every
process condition was different and adjusted to deposit thin ﬁlms of at least 3 um in
thickness allowing for accurate composition measurement in EDS analysis. Also, a graph
of composition as a function of plasma voltage is shown in Figm'e 4-11. It shows that the
deposited thin films had higher Ni concentration than the target alloy and the Ni
concentration increased with increasing energy. , I

Unlike the target erosion test, the thin ﬁlm samples in this test and the next three
tests were all deposited after a steady-state was attained, following the experimental
procedure described in section 3.3. In the present experiment, since the processes were
controlled by voltage regulation mode, the changes of exponent n can alter the discharge
current during deposition (Equation (4.6)). The changes of discharge current may
inﬂuence the target surface topography and change the sputtering yields. However, the
change rate of discharge current is slow and a steady-state condition can be easily
recovered during every change. Thus, it is reasonable to assume that a steady-state
condition was maintained for every process condition during deposition, i.e., the
sputtered ﬂux leaving the target surface always carried the same composition as the target
alloy.

85

 

 

 

 

 

 

500 V 0.53 Pa, 35—>38 W, 63—)70 mA. time=150 min, D=51 mm
Position Center
Thickness 31394 A
Dep. Rate 209 Almin
Composition (Ni at%) 52.79i0.11%
625 V 1.53 Pa, 284112 W, 151—)184 mA, g’m§=12()min, D=§1 m
Position Center
Thickness 77974 A
Dep. Rate 650 Almin
Composition (Ni at%) 56.0210.11%
750 V 0.53 Pa, 128—)162 W, 165-9211 mA, time=90 min, D=51 mm
Position Center
Thickness 74146A
Dep. Rate 824 A/tnin
Composition (Ni at%) 56.12:I:0.07

 

 

Target Composition (Ni-Ti at%) 50.59 - 49.41 at% :I:0.04 Ni

 

 

Table 4—4 Porcess Conditions and the Compositional Results in the Voltage Effect

Test

57
56
55 -
54 '
53 '
52 '
51 -

I

50 '
49
48
47

Composition (Ni at%)

 

Figure 4-11

m
u).

_______________________________ Sputter Target
Composition

625 V

Plasma Voltage

Composition Change as a Function of Plasma Voltage

 

 

86

Since the deposition processes were under steady-state condition and the gas
pressure was kept constant, two other factors are left to be considered as sources for
causing the composition change: the effects of angular distribution of sputtered atoms, and
resputtering. An important issue which must be noted here is that thin ﬁlms deposited in
this test were made after the target erosion test and the same eroded target was used for
these depositions. In section 4.2.1, the macroscopic target surface change led to a
different angular disuibution of sputtered atoms and could be a potential factor causing Ti-
enrichment in thin ﬁlms. If this were true, thin ﬁlms deposited in this test must be Ti-rich.
However, the opposite results were found in this test using the eroded target. This
implies that when compared to other factors, such as resputtering effect and changed
sputtered ﬂux effect, the change of angular distribution of sputtered atoms due to the target
localized erosion effect is not very signiﬁcant.

As discussed in section 4.1.4, the neutral atoms causing resputtering come from
the primary bombarding ions reﬂected from the target surface. Thus, the larger number of
bombarding ions induces more neuual atoms and increases the resputtering effect. Also,
the depleted amount of Ti increases as the resputtering effect increases. Since the number
of bombarding ions can be estimated by the product of the discharge current and the
deposition time, the resputtering effect resulting from the neutral atoms can be then
decided. Referring to Table 4-4, the number of bombarding ions at 625 V was about
twice as many than those at 500 V. Therefore, more Ti was resputtered from the thin ﬁlm
deposited at the 625 V process condition. Under conditions of 625 V and 750 V, the
numbers of bombarding ions were similar. Thus, the resputtering effect was almost the
same, as were the compositional results. The energies of the neutral atoms should
inﬂuence the resputtering rate. However, Chang's study [26] showed the component
resputtering yield ratio Y-r/ YNi had little change when it was above 500 eV. So

basically, the ratio can be assumed to be constant at these energy levels.

87
From this test, it was found that the change of angular distribution of sputtered

atoms was not evident when compared to the other factors. Also, the resputtering effect

can induce a signiﬁcant composition difference in the deposited thin ﬁlms.

4.2.3 Effect of Plasma Power

Theoretically, plasma power is not-thought to be a factor affecting the
compositional shift between thin ﬁlms and target alloy in magnetron sputtering deposition.
However, the magnetron sputter-deposition processes controlled by power regulation
mode have been practically used in the semiconductor manufacturing industry because of
the safety and sputtering rate issues. So, it is worthwhile to verify if processes controlled
by power regulation mode can be found to cause any shift in thin ﬁlm composition. If no
shift is evident, the power regulation mode may be used to control the deposition
processes in a practical manner. I.

Table 4-5 shows the process conditions and the compositional results in the
plasma power effect test. Processes were controlled by power regulation mode and
processed at 150 W, 300 W, and 450 W, respectively. Pressure, 0.44 Pa, and working
distance, 50.8 mm, were always kept constant at the three power levels.

There is a set of data for each of two different positions on the substrate. One is at
the position of target center axis, the other one is at radius 25.4 mm, angle zero degree in
the polar coordinate system relative to the center position. The thickness proﬁle over the
substrate showed a maximum at the center. Maissel [76] suggested that the conditions for
ensuring better uniformity of deposition are obtained when a plane cathode is about twice
as large as hat of the substrate used. Because in the study, thin ﬁlms were deposited by a
50.8 mm sputtering cathode, it is expected that a uniform ﬁlm thickness and composition

cannot be attained over the substrate of the same size as the cathode source.

88

 

 

 

 

 

 

 

 

150 Watts 0.44gk947—t830 v, 153-»175 mA, time=80 min,D= 51 _m__rrt]
Position At Center 25mm Off axis
Thickness 88,959 A 60.98411
Dep. Rate 1112 Almin 762 Almin
Composition (Ni at% 5232100596 51.59:t0.26%
300 Watts 0-44 Psalm->918 VJQEW
Position At Center 25mm off axis
Thickness 78,293 A 56,358 A
Dep. Rate 1957 A/min 1408 A/min
Composition (Ni at% 49.7210.30% 50.64:t0.19%
450 Watts 0.44 Pa, 932—)983 v, 4824465 mA, time=20 min, D=51 mm.
Position At Center 25mm off axis
Thickness 58,756A 56,433 A
Dep. Rate 2938 A/min 2821 A/min
Composition (Ni at% 51.15:t0.12% 51.55:t0.l6%

 

 

 

 

Target Composition (N i-Ti at%) 50.59 - 49.41 at% :I: 0.04 Ni

 

 

Table 4-5 Process Conditions and the Compositional Results in the Power Effect

Test

52.5
52.0

ﬂ
1

51.5 .-
51,0.
50.5L:
50.0 ~

Composition (Ni at%)

49.5 '-

 

49.0 J

—"‘3'_ Center

—-0— 25mm off axis

1 I 1

2:22:22. ::::::....:::::::::::'.::1 SpuuerTarget

Composition

 

150 Watts

300 Watts

Plasma Power

450 Watts

Figure 4-13 Composition Change as a Function of Plasma Power

89

 

 

".7

E 2.5

U

E

o

7‘ 2.0——

F I5

{5 ' _ Ne

2

in Ar

0 l.O--

.J Kr

'3?

S 0'5"" Cu CATHODE

LU - .

'20.;0 I II I I I I I
O I 2 3 4 5, 6 7

DISCHARGE CURRENT (A)

Figure 4-12 Local Gas Density Measured Near the Magnetron Cathode During
the Sputtering of Cu in Ne, Ar, and Kr as a Function of Discharge Current [78]

 

90
Note that at 450 W, the voltage increased during deposition. This result was

opposite from that obtained at conditions of 150 W and 300 W in this test and at 250 W in
the target erosion test, whose voltage increased during deposition. One possible reason
causing this abnormal phenomenon at 450 W is the rarefaction effect. One characteristic
of magnetron plasmas is a high level of sputtering of the cathode. Sputtered atoms emitted
from the bombarded surface can have gas-phase collisions with background gas atoms
and, as such, transfer some of their energy to the atoms through elastic collisions. The
effect is known as thermalization. During this process, signiﬁcant amounts of thermal
energy are transferred to the gas atoms, and their temperature can rise signiﬁcantly [77,
78]. Because of the higher atom velocities near the cathode, the gas may become rareﬁed
ﬁeld, and the local density near the cathode can decrease. This is shown in Figure 4-12
[7 8]. This effect is strongly dependent on the thermal conductivity of the gas, the sputter
yield of the target materials, and the applied discharge power. Ions in the plasma are
created by ionization of the backgron gas. Driving down the local gas density by
sputtered atom heating results in a more resistive plasma. Thus, the local gas density near
the cathode decreases with increasing discharge current. This implies the creation of a
more resistive plasma, i.e., the decrease of exponent n in Equation (4.6). This may cause
the voltage increase dming deposition when the process is controlled by a high power, as
at 450 W in this test.

A graph of composition as a function of plasma power is shown in Figure 4-13. It
shows that at 300 W and the center position, a Ti-rich (with respect to the target alloy) thin
ﬁhn could be produced; while at 150 W and 450 W, Ni was always the major component
in the deposited ﬁlms.

As discussed in section 4.1.4, resputtering effect results from the bombarding of
neutral atoms, which are the primary bombarding ions reﬂected from the target surface.
Since the center position has a greater chance of being bombarded by the reﬂected neutral

atoms than the off center position, thin ﬁlms deposited at the former position will have

91
higher Ni concentrations than those deposited at the latter position. However, in addition
to the neutral atoms, the other resputtering source is the ion ﬂux introduced by the
"unbalanc " magnetron. Referring to Figure 4-8, the ion ﬂux from the central is less
than that of the outer region. Thus, the outer region will suffer more ion ﬂux
bombardment and cause a greater amount of Ti depletion from the deposited ﬁlms than the
center position, and this is opposite to that caused by neutral atoms. The resputtering ion
ﬂux is very dependent on discharge current. Thus, the discharge crurent will determine
which resputtering source dominates during the deposition process. The experimental
results from this test and the next test (section 4.2.4) showed that when the discharge
current was above about 250 mA, the outer region of the substrate would sustain more
bombardment of ion ﬂux and cause more loss of Ti component in the thin ﬁlms.

Since the processes in the present test were controlled by power regulation mode,
the arguments about the non-steady—state sputtered ﬂux discussed in section 4.2.1 could
be applied to the present test, i.e., if the voltage continuously decreased during deposition,
the sputtered ﬂux leaving the target surface had more Ti component than the target alloy,
and vice versa. At 450W, the voltage abnormally increased during the deposition due to
the rarefaction effect and resulted in the sputtered ﬂux having more Ni component than the
target alloy. Upon combining the resputtering effects, the deposited ﬁlm composition
possessed higher Ni concentration than the target alloy.

Referring to Table 4-5 and comparing the process conditions at 150W and 300W
at the center position, both of them had very similar process conditions: the amount of
decreasing-voltage during deposition, the average plasma voltage, and the number of
bombarding ions. Through the discussions of the above two tests, the compositions of
the deposited ﬁlms ﬁom these two process conditions were predicted to be very close.
However, the experimental composition resulting ﬁom each condition was signiﬁcantly
different. Note that the only large difference between the two process conditions was the

discharge current. As discussed in target erosion test, the sputtered ﬂux carrying more Ti

92
component than the target alloy is due to the existence of transition stages before the re-

establishment of steady-state. And the total amount of the excess Ti is determined by the
mechanisms occurring dming the transition stages. For a multiphase material under high
ﬂuence sputtering, the topography effect is the dominant factor during the transition
periods, i.e., steady-state is reached when the rates of cone formation and annihilation are
equal. Unfortunately, the processes governing the evolution of surface topographies are
still under discussion and inconclusive. The theoretical sputtering models for such
systems are also not available. Thus, it is difﬁcult to explain the detailed processes
happening during the transition periods and its effect on the sputtering yields.

However, based on the process conditions of 150 W and 300 W in this test, one
possible reason for the compositional differences between them is that at 300 W, the
higher level of discharge cmrent (~ 312 mA) resulted in higher density of cones on the
target surface than those at 150 W (~ 166 mA), as described in section 4.1.3. Because the
transition periods were determined by topography effect, it would take a longer time for
the 300 W setting to establish the steady-state condition than it would the 150 W setting.
Therefore, the sputtered ﬂux at 300 W had more possibilities for a Ti-rich situation than
that at the 150 W setting. The ﬁlm deposited at 300 W thus had a higher Ti concentration
than that deposited at 150 W. But, this is only a qualitative description of the comparison
between the two process conditions. For a quantitative understanding of this effect,
further experiments under deﬁned conditions (Ultra High Vacuum, mass analyzed ion
beam) are needed to obtain more information and insight into the development of the

surface features, and the relationship with the sputtering yields.

93
4.2.4 Effect of Gas Pressure

After ejection ﬁom the target surface, the sputtered atoms are transported to the
substrate in the deposition chamber. It is likely that the sputtered atoms can have
collisions with background gas atoms. Also, the reﬂected neutral atoms responsible for
the resputtering of growing ﬁlms can suffer collisions with the gas atoms before reaching
the substrate. Thus, the gas pressru'e plays an important role during transport for the
sputter-deposition method. In this test, processes were controlled by voltage regulation
mode. Thin ﬁlms were deposited at 0.44 Pa, 0.64 Pa, and 1.01 Pa, respectively, while
keeping the voltage at 600 V. Table 4-6 shows the process conditions and the
compositional results in the pressure effect test. A graph of composition as a function of
gas pressure is also shown in Figure 414. Since the processes were controlled by
voltage regulation mode, it could be assumed the steady-state condition was maintained
during deposition. Note that the phenomenon of non-uniform thickness and composition
of thin ﬁlms at different substrate positions can be explained by the arguments discussed
in the above test.

However, unlike the plasma power effect test, the trend of changes of composition
with pressure was different at the different substrate positions: at the off center position,
Ni concentration in thin ﬁlms linearly increased with gas pressure; while at the center
position, a minimum Ni concentration could be found at the middle range of gas pressure
(0.64 Pa). From the above test, the major cause resulting in the depletion of Ti at the off
center position was the resputtering effect caused by the ion ﬂux, which was introduced
by the "unbalanced" magnetron. Since the resputtering ion ﬂux is highly dependent on
discharge current and roughly independent of gas pressure (section 4.1.4), the amount of
depleted Ti increased with increasing discharge current. Because the discharge current
increased with increasing pressure, the Ni concentration increased with increasing

pressure at the off center position.

94

 

 

 

 

 

 

 

 

 

 

 

 

0.44 Pa V. 68-988WWJ - =
Position At Center 25mm off axis
Thickness 69,335 A 55239A
Dep. Rate 513 Almin 409 Almin
Composition (Ni at%) 51.76:I:0.05% 51.05:|:0.23%
0.64 Pa 600 V, 159—)195 Watts, 260-)320 mA,time=80 min,D=51 mm
Position At Center 25mm off '5
Thickness 86,411 A 69,748
Dep. Rate 1080 A/min 872 A/min
Composition Qii at%) 50.66i0.17% 51.68i0.28%
1.01Pa muwammsammum
Position At Center 25mm off '5
Thickness 129,700 A 91,859 A“
Dep. Rate 1621 A/min 1148 A/min
Composition (Ni at%) 524110.189!) 52.59i0.25%
Target Composition (N i-Ti at%) 50.59 - 49.41 at% $0.04 Ni

 

 

Table 46 Process Conditions and the Compositional Results in the Pressure Effect

 

 

 

Test

530 P '—"="— Center
52'5 -—°-—25mm off axis

8:0: 52.0

a _

a 51.5

t:

.o

.5 51.0

8.

g 50.5 -------------------------------

U \ Sputter Target
50.0 ,_ Composition
495 = I I : I I I I :

0.44 Pa 0.64 pa 1.01 Pa
Pressure

Figure 4-14 Composition Change as a Function of the Gas Pressure

95
At the center position, thin ﬁlm cornpositions are determined by the deposition of
sputtered atoms and the resputtering of neutral atoms. The gas pressure in the deposition
chamber decides their collision numbers before reaching the substrate. More precisely,
the parameter determining the number of collisions is the mean free path. From the mean

free path equation:
1

«Endldzn
where n is the gas concentration, d, and d, are the molecular diameters of colliding atoms,

A, -.—. (4.8)

one of them being the background gas atom's.

In this study, the sputter gas was Ar, with an atomic diameter of 3.67x10' cm,
and the sputtered atoms were Ti and Ni, with diameters of 1.46x10' cm and 1.25x10'
cm, respectively. Thus, the mean ﬁee path for Ar, Ti, and Ni at different pressures are
shown in Table 4-7.

Table 4-7 The Mean Free Path of Ar, Ti, and Ni at Pressure of 0.44 Pa, 0.64 Pa, and

 

 

 

 

 

1.01 Pa.
0.44 Pa 0.64 Pa 1.01 Pa
1A: (mm 14 10 6
4n (m) 36 26 16
Ma (mm) 42 29 18

 

 

 

 

 

Considering the scattering effect of the sputtered atoms, Ti atoms have a shorter
mean free path than Ni atoms at the same gas pressure. Also, Ti atoms are lighter than Ni

atoms (Table 4-1). Thus, there are more signiﬁcant scattering effects of Ti atoms during

96

the transport process than Ni atoms. Ti atoms are more easily scattered to other places
inside the chamber as well as the substrate. Therefore, the Ti concentration in the thin
ﬁlms will decrease with increasing gas pressure. However, the distance between target
and substrate was 51 mm, resulting in only slight differences in collision numbers
between the sputtered Ti and Ni atoms at the three pressure levels. Therefore, it is
questionable that such slight differences in collision numbers can result in the
compositional changes shown in Figure 4-14. Also, the scattering effects of the sputtered
atoms cannot explain the increasing of Ti concentration when the pressure was increased
from 0.44 Pa to 0.64 Pa. /

The other factor to consider is the resputtering effect by the neutral atoms. When
the pressrn'e increased from 0.44 Pa to 1.01 Pa, the collision number changed from 3 to 8
times. The increasing collision number would reduce the resputtering effect. Therefore,
the Ti concentration in the thin ﬁlms will increase with increasing gas pressure. But, the
increasing of collision numbers can also increase the effect of thermalization and the gas
becomes rareﬁeld near the cathode, as discussed in section 4.2.3. This rarefaction effect
changes the local number of gas atoms that can scatter the sputtered atoms and the neutral
atoms. Thus, the increasing pressure (up to 1.0 Pa) increases the rarefaction effect and
increases the resputtering effect again. This phenomenon causes more depletion of Ti.
So, a maximum Ti (minimum Ni) concentration can be found when increasing from low
pressure (0.44 Pa) to high pressure (1.01 Pa). This argument can also explain why at a
high pressure region (1.01 Pa) the ﬁhn deposited at the center position had a comparable
composition to that deposited at the off center position, which suffered the sputtering of
ion flux.

This test indicates that at pressures lower than 1 Pa, the resputtering effects will
dominate the scattering effect of sputtered atoms and determine the ﬁnal composition

results of the TiN i alloys.

97
4.2.5 Effect of Working Distance

Changing the distance between target and substrate has the same effect of changing
gas pressure — both methods change the number of collisions of the sputtered atoms and
reﬂected neutral atoms before they arrive at the substrate. In this test, due to the limitation
of the sputtering system conﬁguration, there were two different working distances, 51
mm and 76 mm, being surveyed. Table 4-8 displays the detailed process conditions and
their compositional results. The processes were controlled by voltage regulation mode
and kept constant at 600 V, and processed at constant pressure, 0.61 Pa. A graph of
composition as a function of working distance is shown in Figure 4-15. Note that the
composition in thin ﬁlms deposited at the two different distances was almost the same
when considering the effect of measurement error.

From the mean free path equation, Equation (4.8), at pressure 0.61 Pa, the mean
free paths for Ar, Ti, and Ni are: AA, = 10 mm, M, = 26 mm, and AN; = 30 mm. As
discussed in above test, the collisions of the neutral atoms determined the compositional
results. When working distance increased from 51 mm to 76 mm, the collision number of
Ar atoms changed from 5 to 7 times. In this collision region, the change of collision
number increased the thermalization effect, and indirectly caused a higher resputtering
effect. Therefore, the Ni concentration increased as the working distance increased.
However, because the increasing resputtering effect was due to the decrease of local gas
density near the cathode, as the working distance increased, this localized span no longer
dominated the transport distance. Thus, the total resputtering effect between these two
working distances showed no signiﬁcant difference, nor did the thin ﬁlm composition.

Due to the limitation of the number of samples, the effect of working distance on
TiNi thin ﬁlm composition cannot be fully exploited. However, in principle, changing the
working distance has the same effect as changing the gas pressure. So, the compositional
change in thin ﬁlms due to changing working distance can be explained by the gas

pressure effect study.

98

 

Distance = 51 mm 600 V, 0.61 Pa 24l—>262W 399—)431 mA time=35 min

 

 

 

 

 

Position At Center Position
Thickness 53,478 A
Deposition Rate 1528 A/min
Composition (Ni at% 51.47:I:0.33%
Distance = 76 mm , 600 V. 0.61 Pa, 281—)286 W, 467-)474 mA, time=55 min
Position At Center Position
Thickness 45,376 A
Deposition Rate 825 Almin
Composition (Ni at% 51.79zt0.05%

 

 

Target Composition (N i-Ti at%) 50.59 - 49.41 at% 10.04 Ni

 

Table 4-8 Process Conditions and the Compositional Results in the Working Distance
Effect Test

 

 

 

51.8 v ,
51.6 9- j
2.3 51.4 ..
z 51 2 "
a o
a .
g 51.0 Sputter Target
g 50 8 -- Composition
0 .
50.6 ‘EIIZZZZZIZIZZZZIZZZZZZZIZZIIIZZZIIZZZZé'.
50.4 1 =
51 mm 76 mm
Distance

Figure 4-15 Composition Change as a Function of Working Distance

 

CHAPTER 5 CONCLUSIONS

The purpose of this study is to exploit the process conditions which can fabricate
thin ﬁlms with controllable compositions deviating either positively or negatively from the
target alloy composition in magnetron sputter-deposition method. In previous studies, only
Ni-rich thin ﬁlms can be made when using single TiN i target. In the present study, Ti-rich
thin ﬁlms were successfully fabricated by means of controlling the magnetron sputtering
process parameters.

In dc magnetron sputtering, the deposition processes are usually operated in argon
at a pressure of 0.15 to 1.5 Pa and current densities from 4 to 60 mA/cmz. Under these
process conditions, the resputtering of growing thin ﬁlms by the reﬂected neutral atoms is
inevitable and has signiﬁcant effects on the ﬁlm composition. If T'iN i thin ﬁlms were
deposited under steady-state condition, the sputtered ﬂux leaving the target surface with the
same composition as the target alloy, the depletion of Ti can be always found in the thin
ﬁlms due to the preferential resputtering of Ti, no matter what the effects of angular
distribution and scattering of the sputtered atoms. However, in non-steady-state transition
stages, there is a chance that the sputtered ﬂux may carry more T1 component than Ni as
compared to that seen under steady-state condition. Thus, a Ti-rich thin ﬁlm (with respect
to the target composition) deposited during the transition periods can be fabricated when the
amount of excess Ti can overcome the preferential resputtering.

In the present experiments, the power supply of the sputtering system can be
controlled by three regulation modes: voltage, power, or current. When the deposition
processes are controlled by voltage regulation mode, steady-state condition can be
maintained during deposition. However, when the processes are controlled by power
regulation mode, non-steady-state conditions can be setup due to the increasing strength of

magnetic ﬁelds in the TORUS®-2C magnetron source during deposition. Finally, the ion

100
current densities can affect the target sm'face topographies. which in two signiﬁcantly affect

the mechanisms occurring during the transition stages.

Following are listed conclusions of the present study.

(1) A Ti-rich thin ﬁlm can only be fabricated in the processes controlled by the
power regulation mode using the TORUS®-2C magnetron source due to the non- steady-
state conditions.

(2) When the process conditions are at 250 W to 350 W, Ti-rich thin films can be
fabricated.

(3) The angular distribution and scattering effects of sputtered atoms are possible
factors inﬂuencing the deposited thin ﬁlm composition. However, in plasma magnetron
sputtering system, they become insigniﬁcant and negligible in the analysis when compared
to the resputterin g effect and the non-steady-state sputtered ﬂux effect.

(4) Under steady-state, Ni concentration in thin ﬁlms will increase with increasing
voltage due to the increasing effect of resputtering.

(5) The change of pressure and working distance alter the signiﬁcance of the
resputtering effect. When thin ﬁlms are deposited under steady-state condition, the
adjustment of pressure and working distance can only reduce the amount of depleted Ti,
but never produce a Ti-rich thin ﬁlm.

(6) At a pressure of 0.64 Pa, the resputtering effect can be reduced to the minimum
and a thin ﬁlm composition close to the target alloy composition can be made.

BIBLIOGRAPHY

10.

BIBLIOGRAPHY

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