Diffusion Behavior and Phase Formation for Ion Implanted Austenitic

Diffusion Behavior and Phase Formationfor Ion Implanted Austenitic Metal Alloys

Von der Fakultat fur Physik und Geowissenschaften

der Universitat Leipzig

genehmigte

D I S S E R T A T I O N

zur Erlangung des akademischen Grades

doctor rerum naturalium

Dr. rer. nat.,

vorgelegt

von Dipl.-Phys. Johanna Lutz

geboren am 16. April 1982 in Hof

Gutachter: Prof. Dr. Dr. Bernd Rauschenbach

Prof. Dr. Thierry Czerwiec

Tag der Verleihung 16. August 2010

Bibliographische Beschreibung

Lutz, JohannaDiffusion Behavior and Phase Formation for Ion Implanted Austenitic Metal AlloysUniversitat Leipzig, Dissertation101 S., 160 Lit., 36 Abb., 1 Tab.

Referat

Plasma-Immersions-Ionenimplantation (PIII) ist ein Ionenimplantationsverfahren, bei dem

Ionen in einem Energiebereich von 1 - 30 keV genutzt werden. Das zu behandelnde Sub-

strat wird in ein Plasma getaucht. Anschließend werden durch Anlegen negativer Hochspan-

nungspulse positive Ionen aus dem Plasma zur gesamten Substratoberflache beschleunigt und

implantiert.

FeCrNi (austenitischer Edelstahl) und CoCr Legierungen haben in der Medizin einen großen

Anwendungsbereich als metallische Implantate. Eine Verbesserung ihrer Oberflacheneigen-

schaften ist in vielerlei Hinsicht von Interesse, da Legierungsbestandteile uber Korrosion oder

uber Abriebpartikel freigesetzt werden und Nebenwirkungen im umliegenden Gewebe verur-

sachen konnen.

Die Arbeit befasst sich mit der Stickstoffimplantation von austenitischem Edelstahl und CoCr

mittels PIII. Neben Charakterisierung des Diffusionsverhaltens von Stickstoff in den Legierun-

gen und der Phasenbildung wurden mechanische, tribologische sowie elektrochemische Eigen-

schaften untersucht und zwischen beiden Legierungssorten verglichen.

Der Transport des eingebrachten Stickstoffs wird in beiden Legierungen durch Zwischengit-

terplatze realisiert, wobei ahnliche Aktivierungsenergien gefunden wurden. Wahrend des

Implantationsprozesses haben zusatzliche Oberflacheneffekte einen wesentlichen Einfluss auf

die Menge des eingebrachten Stickstoffs. Durch Separation der einzelnen Wechselwirkungen,

die durch Gasmolekule, Ionen aus dem Plasma sowie energetische Ionen entstehen, konnte ein

qualitatives Modell erstellt werden, das die Prozesse an der Oberflache bis hin zur Diffusion

im Volumenmaterial voneinander trennt und beschreibt.

Durch das Einbringen von Stickstoff in das ursprungliche fcc Kristallgitter kommt es in bei-

den Legierungen zu Gitteraufweitungen und zur Bildung der sogenannten γN-Phase. Bei

PIII Prozesstemperaturen uber 400 ◦C kommt es zur Prazipitation von CrxN, verbunden mit

einer an Chrom verarmten Matrix. Wahrend diese fur austenitischen Edelstahl eine ferriti-

sche (bcc) Struktur bildet, findet man bei CoCr Legierungen eine fcc CoNi Phase, die jeweils

auf unterschiedliche Weise das Diffusionsverhalten beeinflussen. Es konnte gezeigt werden,

dass die Ausscheidung von CrxN nicht nur temperatur- sondern auch zeitabhangig ist. Ein

Zwei-Schichten-Modell bestehend aus CrxN/γN/γ-Bulk mit unterschiedlichen Schichtwachs-

tumsgeschwindigkeiten wurde daraus abgeleitet.

Die behandelten CoCr und Edelstahl Substrate weisen sowohl eine gesteigerte Harte als

auch eine verbesserte Verschleißresistenz auf. Die elektrochemischen Eigenschaften der CoCr

Legierungen sind jedoch deutlich beeintrachtigt. Wahrend bei Edelstahl erst ab Prozesstem-

peraturen uber 400 ◦C eine Verminderung der Korrosionsresistenz eintritt, findet man eine

erhohte selektive Abgabe von Kobaltionen nach PIII bei allen Temperaturen. Das Einbrin-

gen von Stickstoff schwacht die Stabilitat des Co-Cr-Ni Verbunds und isoliert dadurch die

Kobaltatome, die so einfacher herausgelost werden konnen.

Contents

1 Introduction 7

2 Fundamentals 112.1 Plasma Physical Basics of Plasma Immersion Ion Implantation . . . . . 11

2.2 Ion-Surface Interaction . . . . . . . . . . . . . . . . . . . . . . . . . . . 14

2.3 Diffusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

3 Materials and Experimental Methods 193.1 Austenitic Stainless Steel and CoCr Alloys . . . . . . . . . . . . . . . . 19

3.2 Plasma Immersion Ion Implantation Setup . . . . . . . . . . . . . . . . 23

3.3 Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24

3.3.1 Secondary Ion Mass Spectroscopy (SIMS) . . . . . . . . . . . . 24

3.3.2 Glow Discharge Optical Emission Spectroscopy (GDOS) . . . . 25

3.3.3 Inductively Coupled Plasma Optical Emission Spectroscopy (ICP-OES) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

3.3.4 X-Ray Diffraction Analysis (XRD) . . . . . . . . . . . . . . . . 27

3.3.5 Scanning Electron Microscopy (SEM) and Energy Dispersive X-Ray Analysis (EDX) . . . . . . . . . . . . . . . . . . . . . . . . 28

3.3.6 Hardness and Wear Measurements . . . . . . . . . . . . . . . . 29

3.3.7 Electrochemical Impedance Spectroscopy (EIS) . . . . . . . . . 29

4 Characterization of Nitrogen Diffusion 314.1 Depth Profiling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32

4.2 Temperature and Temporal Dependency of the Layer Thickness . . . . 36

4.3 Summary of Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

5 Influence of Gas and Plasma on Adsorption, Incorporation and DesorptionProcesses 415.1 Retained and Incident Fluence . . . . . . . . . . . . . . . . . . . . . . . 42

5.2 Influence of Background Pressure . . . . . . . . . . . . . . . . . . . . . 44

5.3 Comparison of Gas Nitriding (GN), Plasma Nitriding (PN) and PIIINitriding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

5.4 The Role of a Surface Oxide Layer . . . . . . . . . . . . . . . . . . . . 48

5.5 Summary of Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52

5

Contents

6 Phase Formation 556.1 Lattice Expansion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 556.2 Temporal Decomposition of Expanded Austenite . . . . . . . . . . . . . 58

6.2.1 Phase Formation after PIII during Different Annealing Times . 586.2.2 Phase Formation after PIII with Different Process Times . . . . 61

6.3 Lattice Parameter and Nitrogen Content . . . . . . . . . . . . . . . . . 636.4 Summary of Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . 65

7 Surface Properties 677.1 Corrosion Behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 687.2 Nano Indentation and Wear . . . . . . . . . . . . . . . . . . . . . . . . 707.3 Tribocorrosion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 747.4 Summary of Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

8 Summary and Conclusions 81

Bibliography 85

List of Figures 95

List of Tables 97

Acknowledgements 99

Curriculum Vitae 100

Selbststandigkeitserklarung 101

6

1 Introduction

Metals are extremely useful as components in medical devices like endoprostheses,

fixation of fractures and intravascular implants. Besides Ti alloys, CoCr alloys and

austenitic stainless steel are the most common metals used in orthopedics [1] due to

their beneficial combination of mechanical strength, ductility and biocompatibility [2].

They are especially used for load-bearing components like total hip replacements, where

ceramics still show a non-negligible risk of sudden brittle failure, whereas regenerative

tissue is still too soft. In spite of these excellent properties, these metals are not without

risk of adverse effects. Especially, the generation of nanoparticles by mechanical wear

processes [3] and electrochemical corrosion in combination with toxic ion release such

as cobalt, chromium and nickel [4, 5] are responsible for implant failure, inflammations

and fatigue fractures [6].

Conventional alloying and metallurgical processing has a long tradition for improving

or adjusting the bulk porperties by inserting additional chemical elements or chang-

ing the microstructure. Besides yield strength, fatigue strength and elastic modulus,

the corrosion resistance is of major interest for biomedical applications. However, a

distinction between bulk and surface which interacts with biological tissue has been

established in the last decades. Thus, the focus is on the development of surface treat-

ments which are able to improve the tribological properties and to prevent or reduce

the release of potentially harmful metal ions. Additionally, CoCr alloys and austenitic

stainless steel also find further application in current and future industrial fields such

as in chemical, aeronautic and automotive industry [7]. Likewise, an improvement of

the surface properties is often desirable in those fields of application, too.

Surface modification by inserting nitrogen ions with plasma and ion implantation tech-

niques offers the possibility of tailoring and improving the properties of metallic materi-

als. These techniques comprise plasma-assisted nitriding, low energy ion implantation,

plasma immersion ion implantation (PIII) as well as conventional beam line ion im-

plantation. The main differences between these methods are the varying ion energy

and the relative fraction of energetic ions, electrons, thermal atoms and ions impinging

on the surface [8].

7

1. INTRODUCTION

The focus of this work will be on plasma immersion ion implantation, a combination

between conventional ion implantation and plasma processing [9–11], where a continu-

ous low pressure plasma (p ≤ 1 Pa) is used to generate ions, which are extracted by

applying short high voltage pulses to the substrate and subsequently implanted into the

sample. PIII is a very attractive technique for industrial application as it circumvents

the line-of-sight limitations of conventional beam implantation and complex-shaped

surfaces can be implanted simultaneously. In addition, the rather high ion current

densities at moderate temperatures (∼ 400 ◦C) lead to thick nitrogen enriched surface

layers within moderate process times of some hours.

It is well known that austenitic stainless steel which is rather soft and prone to ex-

cessive wear rates, can be surface hardened by nitriding. A characteristic expanded

lattice (designated as expanded austenite, S-phase, m-phase or γN [12–14]), at process

temperatures up to 400 ◦C is formed which is accompanied by improved mechanical

and tribological properties [15, 16]. In contrast, only a few publications exist about

surface modification with nitrogen of CoCr alloys where a similar lattice expansion

is observed [17–21]. However, no comprehensive investigation of the influence of the

alloys’ composition, ion implantation temperature and further process conditions on

diffusion and phase formation as well as on the resulting surface properties has been

published until now. In the cited references, only one or two alloys at a restricted set

of treatment parameters were investigated.

The aim of this work is a detailed investigation and comparison of nitrogen PIII into

austenitic Fe-Ni-Cr-Co alloys, namely CoCr alloys and austenitic stainless steel. Be-

sides, a systematic study of nitrogen diffusion at different process conditions, the phase

formation and the resulting surface properties will be investigated. The influence of

the alloy composition and the process parameters will be studied.

After giving an insight into plasma immersion ion implantation and the respective

plasma- and ion-surface interactions present during implantation (chapter 2), a short

overview about the used analytical techniques will be presented (chapter 3). The first

main part is dedicated to nitrogen diffusivity (chapter 4) and the resulting layer for-

mation in CoCr and stainless steel. The diffusion behavior at different PIII process

conditions will be investigated focusing on different process temperatures and times.

It will be shown that the incorporation of nitrogen during PIII is not only influenced

by implantation but also by further processes taking place during and after the high

voltage pulses.

To understand their basic mechanisms and the contribution of gas and plasma to

the layer formation, experiments separating the different surface processes with their

8

respective active species from each other were performed and will be analyzed in chap-

ter 5. The influence of nitrogen background pressure and of an oxide layer will be

presented. Again, both types of alloys - CoCr and austenitic stainless steel - will be

compared to highlight similarities and differences for both materials.

As diffusion is strongly correlated with the atomic lattice structure and precipitation

inside the alloys, a detailed investigation of the phase formation is indispensable. In

chapter 6 the stability of the expanded austenitic phase and the influence of temper-

ature and time on its decomposition will be examined. Additionally, the focus will

further be on the nitrogen content in correlation with possible atomic nitrogen sites in

the bulk material and the resulting lattice expansion.

Finally, at the end, the surface properties of the PIII nitrided CoCr alloys will be

characterized in regard to their mechanical, tribological and electrochemical behavior

(chapter 7). Especially, the wear mechanisms in physiological solution will be studied

considering the application of CoCr alloys for medical implants. For austenitic stainless

steel, ample results can already be found in literature.

9

2 Fundamentals

Plasma immersion ion implantation is a versatile method for surface modification which

was developed independently in the United States and Australia about 20 years ago

[9–11]. Negative high voltage pulses are applied to a substrate which is immersed in a

plasma. Thus, positive ions are extracted from the plasma and simultaneously acceler-

ated towards the whole surface. PIII is an ideal technology for complex shaped surfaces

since it circumvents the line-of-sight restrictions of conventional ion beam implanters.

Hence, large complex areas can be surface modified without beam steering in a short

and cost-efficient way.

In the following chapter, the basic scientific principles behind plasma immersion ion

implantation are presented. At the beginning, some basics of plasma physics are in-

troduced, especially the theoretical description of macroscopic plasma wall interaction

with time and varying voltages. A second part is dedicated to microscopic ion-surface

interactions since the comprehension thereof is a basic condition to understand the

subsequent layer formation.

2.1 Plasma Physical Basics of Plasma Immersion Ion

Implantation

Plasma Parameters

A plasma is a quasi-neutral gas of charged and neutral particles which exhibits collective

behavior [22]. Quasi-neutrality describes the apparent charge neutrality of a plasma in

whole, whereas at smaller length scales, the positive and negative charges in the plasma

may give rise to locally charged regions and electric fields. Collective behavior describes

the ability to shield electric potentials. A potential Φ0 in a plasma at the point x = 0

[23] can be described using Poisson’s equation. Thus, the resulting screening or decay

11

2. FUNDAMENTALS

of the potential Φ is

Φ = Φ0 exp

(−|x|λD

)(2.1)

by displacing positive and negative charges near the external potential. λD is the

so-called Debye length

λD =

√ε0kBTee2ne

, (2.2)

with the dielectric constant ε0, the Boltzmann constant kB, the electron temperature

Te, the plasma density ne and the elementary charge e. The Debye length λD is a

measure of the screening length inside the plasma, forming a space charge region where

charge neutrality is not maintained. With increasing density λD decreases since more

electrons are available for shielding. At the same time, a higher electron temperature

increases λD as a higher mobility is encountered. Thus, one criterion for an ionized gas

to be a plasma is that λD << L, where L is the dimension of the system where the

plasma exists. In this case an externally imposed potential in the plasma is shielded

within a short distance, thus keeping quasi-neutrality of the whole plasma. For typical

laboratory plasmas, the Debye length is small. For an electron temperature Te of 4 eV

and density ne of 1010 cm−3 a Debye length of about 140 µm is found.

However, a precondition for shielding is the existence of enough particles to form the

shielding cloud. The number of particles in the so-called Debye sphere with radius λD

is

ND = ne4

3πλ3D. (2.3)

Hence, in addition to λD << L, the number of particles has to be much greater than

one (ND >> 1) to allow collective behavior.

A third condition arises due to collisions in the plasma and is characterized by the

plasma frequency ωpe which describes oscillations of charged particles. Since Coulomb

interaction has to be the dominant interaction mechanism in the plasma as compared

to hydrodynamic interactions (collisions with neutral particles) ωτ > 1 is required [22].

τ is the mean time between collisions with neutral atoms.

Static Plasma Sheath

A substrate which is immersed in a plasma will be negatively charged compared to

the plasma due to the higher mobility of the light electrons in respect to the ions.

The repulsion of electrons results from the formation of a positive space charge region

shielding the neutral plasma from the negative substrate. The typical extension of the

12

2.1. PLASMA PHYSICAL BASICS OF PLASMA IMMERSION ION IMPLANTATION

plasma sheath xs is given by the electron Debye length λD. At the same time, ions are

accelerated to the surface through this sheath. However, to maintain the screening the

so-called Bohm criterion has to be fulfilled [24]. This condition for sheath formation

demands that the ion speed at the sheath edge must be equal or greater than the ion

sound velocity uB = (kBTe/mi)1/2 [25] (mi is the ion mass). Therefore, the plasma

wall transition region is divided further into a positively charged sheath and into a

quasi-neutral presheath with a small potential drop Φ ≈ kBTe/2e and a reduced ion

density towards the sheath edge n(xs) = ni exp−1/2 ≈ 0.6 ni, compared to the uniform

plasma with density n0. Consequently, the ion drift current density which enters the

plasma sheath from the presheath is

Ji = exp−1/2 e n0 uB ≈ 0.6 e n0 uB. (2.4)

Using Poisson’s equation and a plane geometry equation (2.4) can be rewritten as the

Child-Langmuir law for potentials larger than kBTe/e [26]

JCL = Ji =4ε09x2s

√2e

mi

Φ3/2. (2.5)

Dynamic Plasma Sheath

During plasma immersion ion implantation a sudden negative high voltage pulse is

applied to the substrate, starting in general from the wall at floating potential. At first,

on the time scale of the inverse electron plasma frequency ω−1pe , electrons near the wall

are driven away towards the bulk plasma, leaving a uniform density ion matrix sheath.

Next, on the slower time scale of the inverse ion plasma frequency ω−1pi , ions within the

sheath are accelerated towards the substrate and implanted therein. Consequently, the

ion density inside the sheath region decreases and causes a corresponding additional

expulsion of electrons and therefore expansion of the sheath edge. Hence, the ion drift

current entering the sheath from the presheath (equation (2.4)) has to be supplemented

by an additional term dxs/dt describing the sheath expansion

Ji = 0.6 en0

(dxsdt

+ uB

). (2.6)

This equation is valid for planar geometry of the substrate, additional geometric cor-

rection factors have to be added when using cylindrical or spherical symmetries [27].

On a longer time scale, the system will establish a new equilibrium and another steady-

13

2. FUNDAMENTALS

state Child-Langmuir sheath is obtained for the new potential. Using equation (2.6)

together with the Child-Langmuir law of equation (2.5) the expansion of the plasma

sheath can be calculated according to Lieberman [28]

dxsdt

=4

9

ε0Φ3/2

0.6en0x2s

√2e

mi

− uB. (2.7)

The derivation of equation (2.7) assumes that the voltage and the sheath width do not

change during the transport of the ions through the sheath. In general, this assumption

is valid for typical plasma parameters [28] and good agreements were found between

this model and experiments measuring the sheath expansion [29–31]. In most cases the

final width of the plasma sheath is in the order of a few centimeters. For instance, at

a pulse voltage of -10 kV, a plasma density of 1010 cm−3 and an electron temperature

of 1.1 eV the plasma sheath width is about 5.5 cm while the Debye length reaches

100 µm.

At the end of a voltage pulse the depleted sheath breaks down and is refilled with

plasma, in the ideal case having the same density as before at the start of the next

pulse.

2.2 Ion-Surface Interaction

Ion Implantation

When energetic ions or neutral particles impinge on a solid target, a variety of different

processes will occur due to elastic (nuclear) and inelastic (electronic) collision events:

ejection of particles and radiation from the target, modifications of the structure of the

target and implantation. The possible processes mainly depend on the energy of the

primary ions. Since ion implantation is a well-established subject with a considerable

amount of literature available [32–36, and references therein] only a short overview

considering low energies present at PIII (1 - 50 kV) is given. In this energy range

nuclear stopping, i.e. elastic collisions of the penetrating ion with target atoms, plays

the dominant role and electronic stopping, can be neglected. The ions penetrate into

the target and lose their energy until they are stopped.

In most applications the projected range Rp is of interest. It is defined as the total

path length of the projectile measured along the direction of incidence. A simple

14

2.2. ION-SURFACE INTERACTION

approximation neglecting electronic stopping is given by Lindard et al. [34]

Rp ≈R

1 + M2

3M1

. (2.8)

However, for practical uses simulation programs like TRIM (transport of ions in matter)

[36] based on analytical calculations and Monte Carlo simulations are used.

In contrast to conventional beam implantation, additional particles are present in PIII

and can be co-implanted [37, 38]. These particles either arise from sputtered material

or from interaction of the plasma with the chamber walls. A high amount of secondary

electrons is produced as a secondary electron coefficient up to ten [39] is observed for

energetic ions. These fast electrons impinge on the chamber walls, where they might

create X-rays and additional electrons or ions which lead to an increase of plasma

density [31, 40]. Further mechanisms are charge exchange collisions within the plasma

sheath [41] which may contribute to an enhanced ion current per pulse as well as inter-

action during the pulse pauses. Finally, it can be shown that the energy distribution of

the ions depends on the ion mass mi and the pulse rise time tr and can be determined

to be tr/m1/2i [42]. Thus, for lower ion masses shorter rise times are necessary to obtain

the same energy distribution and therefore the same depth distribution.

Sputtering

Another important process during implantation is sputtering where momentum is

transferred from ions to atoms in the near surface layer, eventually allowing an es-

cape from the surface across the surface energy barrier. The process is characterized

by the so-called sputtering yield Y which is defined as the mean number of emitted

atoms per incident particle. It does not only depend on the composition and structure

of the target but also on species, energy and ion incidence angle of the incoming beam.

According to Sigmund [43] the sputtering yield can be expressed as

Y = ΛFD(E, θ, x = 0). (2.9)

Λ contains the material properties like density and surface binding energies whereas

FD(E, θ, x = 0) is the energy deposited into nuclear processes at the surface and can be

written as FD(E, θ, x = 0) = αNSn(E). N is the atomic density and α describes the

ratio of the target atom mass to incident atom mass, which is an analytical expression

[44]. Using an appropriate nuclear stopping cross section Sn(E), a numerical value for

15

2. FUNDAMENTALS

the sputtering yield can be determined [36].

Sputtering does not only influence the implantation process by eroding the surface

but also limits the concentration of the implanted ions in the target. During sputter-

ing target atoms as well as already implanted atoms are eroded. Therefore, after a

certain time an equilibrium is reached where as many implanted atoms are removed

by sputtering as are added by implantation. Thus, the concentration distribution of

the implanted atoms is influenced. Whereas low ion doses lead to a maximum of the

concentration profile situated at the mean projected range Rp, high ion doses cause

constant concentration profiles (see figure 2.1).

Sputtering plays a decisive role during plasma immersion ion implantation. The rather

H

M

L

Concentr

ation

Surf

ace

Sputteredatoms

Ions

Rp

Amount ofsurfaceeroded

Lowdose

Highdose

Mediumdose

Depth

Rp

Figure 2.1: Schematic development of the concentration profiles of the implanted atoms afterlow dose (L), medium dose (M) and high dose (H) ion implantation (after [32]).

low energies (1 to 50 keV) for PIII, compared to usual implantation methods, result in

high sputter yields near the maximum value since the energy is deposited in collision

cascades in the surface-near regions of the target. Additionally, short ion ranges lead

to a fast build-up of implanted atoms. Therefore, the high dose region is often reached

after a short PIII treatment time of only a few minutes [41]. In addition, sputtering is

essential for nitriding of metals like stainless steel. The sputtering removes the native

oxide layer from the surface at the beginning of the process, promoting a faster diffusion

of nitrogen into the bulk (see section 3.1). Nevertheless, re-oxidation from molecules

of the residual gas may take place leading to changes in the sputtering yield [45, 46].

Finally, the surface roughness can also be increased due to the energetic ion bombard-

ment. Hence, changes in the surface structure and topography may occur [47, 48]. The

16

2.3. DIFFUSION

roughening is related to different sputter yields of the crystallographic orientations and

impurities in the material, but is also influenced by ion species and angles used for

sputtering [32].

2.3 Diffusion

The dominating mechanism during layer formation at plasma immersion ion implanta-

tion is diffusion. Due to the insertion of foreign atoms near the surface, a concentration

gradient arises leading to a transport of atoms from the area of higher concentration

to the area of lower concentration. The implantation process can be described by im-

plantation in a semi-infinite medium with an ion range distribution approximated by

a delta function [45]. The implantation depth of 10 - 100 nm can be neglected since

the diffusion range is almost one order of magnitude higher.

Diffusion can be mathematically described by Fick’s first law [49]

~j(~r, t) = −D∇C(~r, t), (2.10)

where j(~r, t) is the flux of particles, D the diffusion coefficient and C(~r, t) the con-

centration of the diffusing atoms. In general D is a symmetric tensor of rank two.

However, for isotropic media (like cubic crystals) this tensor can be reduced to a scalar

quantity. Combining equation (2.10) with the continuity equation, Fick’s second law

can be derived∂C(~r, t)

∂t= ∇(D∇C(~r, t)). (2.11)

If D is concentration-independent Fick’s second law in equation (2.11) for a one-

dimensional system will reduce to

∂

∂tC(x, t) = D

∂2

∂x2C(x, t). (2.12)

Equation (2.12) is only valid for dilute solutions where diffusion arises from solute-

solvent interactions, and solute-solute interactions are negligible. Thus, for concen-

trated solutions, this law has to be amended taking a concentration-dependent diffusion-

coefficient into account [50].

From a mathematical point of view, equation (2.12) is a second order, linear partial dif-

ferential equation needing initial and boundary conditions to find particular solutions.

Considering a constant surface concentration and zero initial concentration (C(x, 0) = 0

17

2. FUNDAMENTALS

for x > 0 and C(0, t) = C0 = constant for t > 0) the solution is [51]

C(x, t) = C0 erfc

(x√4Dt

), (2.13)

where erfc is the complementary error function. The quantity√

4Dt is called diffusion

length and is a characteristic distance in diffusion related processes.

In addition, the diffusion coefficient may be temperature dependent, as diffusion can

be described as a thermally activated process: In order to move to a new location the

atom must have a certain energy (activation energy Ea) to pass potential barriers in

the crystal. Heat supplies the atom with the energy to exceed this barrier. The relation

of D to the temperature T can be expressed with an Arrhenius equation

D = D0 exp

(− Ea

kBT

), (2.14)

where D0 is a material dependent constant. Plotting lnD versus 1/T (Arrhenius plot),

the activation energy Ea can be determined using the slope of the resulting line in

combination with equation (2.14).

Two basic atomic mechanisms which give rise to diffusion in solids after ion implanta-

tion are interstitial and vacancy diffusion which dominate mainly at low temperatures

(T ≤ 0.25 Tm; Tm - melting temperature). Interstitial diffusion means that an atom

diffuses by jumping from one interstitial site to the next. These solute atoms are gener-

ally smaller than the solvent atoms and they usually occupy octahedral or tetrahedral

sites of the lattice. In contrast, vacancy diffusion is characterized by jumping of an

atom into a neighboring vacancy. Typical values of Ea are approximately 1 eV for

interstitial diffusion and 0.5 eV for vacancy diffusion. Finally, it has to be mentioned

that diffusion can be accelerated along the grain boundaries in polycrystalline materials

with lower activation energies for these “defective” regions [52].

18

3 Materials and Experimental

Methods

To understand and elucidate the basic mechanisms during nitriding of CoCr and stain-

less steel a detailed knowledge about the studied materials and the nitriding process is

necessary. Therefore, to get an overview about the dominant mechanisms an investiga-

tion of the chemical composition, phase formation, mechanical properties, morphology

and structure is essential.

In the following chapter, the investigated materials, CoCr and austenitic stainless steel,

are introduced and characterized whereas a detailed view on the employed plasma im-

mersion ion implantation process is provided in the second part. Finally, with the aim

to characterize the samples from different viewpoints, different analytical techniques

are presented.

3.1 Austenitic Stainless Steel and CoCr Alloys

Properties

Stainless steel consists of iron containing less than 1 wt.% of carbon as well as additions

of a minimum of 13 wt.% chromium. The chromium enables the formation of a thin

chrome-containing oxide layer at the surface which passivates the iron in the bulk and

protects it from corrosion. Several other alloying elements like nickel, molybdenum,

manganese, silicon, tungsten, vanadium, cobalt, niobium and titanium can be added to

give characteristic properties such as strength, toughness or hardenability to the alloy

[53, 54]. Depending on their microstructure stainless steels can be divided into three

different groups: martensitic, ferritic and austenitic stainless steels corresponding to a

body centered tetragonal, body centered cubic and face centered cubic crystal lattice,

respectively, with the chemical constituents being randomly distributed on substitu-

tional sites of the lattice.

19

3. MATERIALS AND EXPERIMENTAL METHODS

The focus of this work is on austenitic stainless steel (see figure 3.1) which is the most

widely used group of stainless steels. Due to its particular good combination of material

- Fe, Co, Cr, Ni, W, ...

Figure 3.1: Fcc lattice of austenitic stainless steel and CoCr alloys. The alloying elements(Fe, Co, Cr, Ni, W, ...) are stochastically distributed on the lattice positions.

properties, like excellent corrosion resistance, great ductility and toughness, austenitic

stainless steel has a broad application range in food, chemical, medical and processing

industries. In order to form austenite it is necessary to add about 8 % nickel to stabilize

the austenitic structure when cooling the steel down to room temperature [55].

A quite similar alloy class to austenitic stainless steel are CoCr alloys which consist of

cobalt and chromium as main elements instead of iron and chromium. According to

their chemical composition and previous heat treatments CoCr alloys usually exist as

fcc alloys (see figure 3.1) or as two-phase alloys consisting of a solid solution hexagonal

phase and an intermetallic tetragonal phase. CoCr alloys also find several applications

in industry.

In this work two austenitic stainless steels, DIN 1.4301 (AISI 304) and 1.4571 (AISI

316Ti), as well as three different fcc CoCr alloys (SY21med, L605 and HS188) were in-

vestigated. The alloy compositions in weight percent are listed in table 3.1. Mirror-like

polished flat samples with diameters from 10 mm to 25 mm were cut from rods and

used as substrates. The original structure of all alloys consists of an fcc polycrystalline

lattice with all alloying elements in solid solution, except for small carbide precipitates

(in the case of CoCr). The grain sizes varied between 10 - 50 µm depending on the

alloy.

Nitriding

Nitriding is a widely applied thermochemial surface treatment during which nitrogen

is induced into steel or other types of alloys. The basic aims are to obtain a higher

surface hardness, improved wear resistance and enhanced fatigue life.

Conventionally applied nitriding techniques in industrial production are gas nitriding,

20

3.1. AUSTENITIC STAINLESS STEEL AND COCR ALLOYS

Alloy Co Ni Cr W Fe Mn Si Mo C La B N Ti

304 9 18 bal. 0.07 0.11316Ti 12 18 bal. 2 2 0.08 0.7L605 bal. 10 20 15 3 1.5 0.4 0.1HS188 bal. 22 22 14 3 1.25 0.35 0.1 0.03 0.015SY21 bal. 0.85 28 0.65 0.7 0.75 6 0.085 0.15

Table 3.1: Chemical compositions of investigated stainless steels (304 and 316Ti) and CoCralloys (L605, HS188, SY21) in wt.%. The balance (bal.) is either provided by Fe or Co,depending on the alloy class.

salt bath nitriding and plasma nitriding. Gas nitriding introduces nitrogen into the

surface by using ammonia which dissociates at the metal surface via a catalytic reac-

tion at a suitable temperature and ammonia gas flow rate. Salt bath nitriding is quite

similar to gas nitriding except that the nitrogen activity comes from liquid salt baths

containing the reactive nitrogen. In principle, these methods improve the mechanical

properties of stainless steel, but will generally cause unwanted deterioration of the cor-

rosion performance as temperatures above 500 ◦C are required to provide a sufficient

high nitrogen activity to enable the nitrogen to diffuse into the surface. Therefore,

plasma and ion implantation technologies operating at moderate temperatures have

been employed to improve the mechanical properties without impairing the corrosion

resistance. These methods use energetic ions leading to a more efficient uptake of ni-

trogen.

Additionally, the passive film on stainless steel presents a diffusion barrier for nitrogen

atoms, thus nitriding is not possible. However, plasma and implantation technolo-

gies circumvent this passive layer impenetrability, as sputtering of the surface is a

fundamental part of these processes [56]. Plasma nitriding processes are based on

glow discharge technologies to introduce nitrogen into the surface. A further nitriding

method involving layer formation is physical vapor deposition (PVD) like pulsed laser

deposition (PLD) [57], plasma assisted nitriding (PAN) [58] and reactive magnetron

sputtering [59]. In contrast, ion implantation techniques enable the controlled inser-

tion of nitrogen atoms by bombardment with energetic ions. In conventional beamline

implantation the ions are produced in a plasma source, extracted, separated according

to their mass, accelerated and implanted into the sample. The range of the ions and

hence the modification depth can be systematically changed from a few nanometers

to several micrometers by varying the ion energy from 10 keV up to 1 MeV [32, 60].

The use of broad beam low-energy (≈ 1 keV), high-current-density (≈ 1 mA/cm2) ion

21


beams for surface modification of stainless steel leads to decisive improvements in tri-

bological behavior as well as in corrosion resistance. The high current densities result

in much thicker layers for the same processing times and temperatures as in normal ion

beam implantation and plasma nitriding [61, 62]. At the same time, the combination

of rather low ion energies with high current densities yields high sputtering rates.

To circumvent the line-of-sight restrictions of ion beam implantation plasma immer-

sion ion implantation was developed [9, 10]. PIII combines the advantages of plasma

nitriding and ion implantation in one method as nitriding of complex shaped surfaces

in a short and cost-effective way is possible.

Several publications about nitriding of austenitic stainless steel can be found in liter-

ature describing a diversity of methods. Nitrogen insertion at low temperatures below

400 ◦C leads to improved tribological behavior without impairing the corrosion resis-

tance due to the formation of the so-called expanded austenite also known as S-phase,

m-phase or γN [12–14]. This phase is characterized by an unusual high nitrogen con-

tent of up to 30 at.% in solid solution and a corresponding lattice expansion between

5 - 15 % resulting in a fast interstitial diffusion of nitrogen [56, 63, 64]. Expanded

austenite is metastable and tends to develop chromium nitrides. In addition, hardness

values up to 17 GPa and an improvement of wear resistance of about 3 - 4 orders of

magnitude have been reported [65–67]. The strengthening mechanism is attributed to

a combination of solid solution hardening by the high interstitial atom content, as well

as the presence of stacking faults in the fcc sublattice of substitional atoms [68]. The

corrosion properties usually remain unaltered, however, they can also be improved, in

particular the pitting potential [69–71]. Nevertheless, the exact mechanisms of phase

formation, lattice expansion and diffusion behavior are not understood in detail and

further investigations are required.

The nitriding behavior of CoCr alloys is by far less investigated than that of austenitic

stainless steel and only few publications exist. Depending on the nitriding method

layer thicknesses up to 30 µm have been reported [17]. Nitriding leads to enhanced

mechanical properties resulting in a reduced wear rate of a factor of 30 and a hardness

increase of a factor of 2 - 3 in comparison with the base material [18, 19]. However,

immersion tests reveal a strongly increased metal release rate after nitrogen treatment

[21]. Although the formation of an expanded lattice similar to austenitic stainless steel

has been found at temperatures around 400 ◦C [20, 21] the diffusion process and the

phase formation have not been studied in detail yet.

22

3.2. PLASMA IMMERSION ION IMPLANTATION SETUP

3.2 Plasma Immersion Ion Implantation Setup

The experiments were performed in a UHV chamber with a volume of 150 l and a base

pressure of 10−6 Pa (see figure 3.2). The plasma was produced using an electron cy-

ECR plasmasource

Gas supply

Pyrometer

High voltagepulse generator

Pumpingsystem

Sampleholder

Figure 3.2: Schematic experimental setup of the UHV chamber.

clotron resonance (ECR) plasma source which was mounted on the top of the chamber

[72, 73]. The working principle of an ECR plasma source is presented in figure 3.3 [74].

Microwave excitation at 2.45 GHz is realized by an antenna housing in a quartz cup.

At the same time a permanent magnetic field B = 87.5 mT is applied, leading to a

resonance zone where the applied microwave frequency ωmw and the electron cyclotron

frequency ωec = eB/me coincide. Thus, the electrons are accelerated in phase and gain

energy in this region. Consequently, they collide with the atoms or molecules of the

supplied gas and cause ionization and plasma sustainment.

For the PIII experiments, the ECR plasma source operating at a power of 150 W

generated a plasma with an electron temperature and plasma density of 1.3 eV and

1.6× 1010cm−3, as determined from Langmuir probe measurements directly above the

substrate holder [75]. At a nitrogen gas flow rate of 150 sccm the resulting pres-

sure during the experiments was 0.53 Pa. Additionally, process pressures of 0.33 and

0.82 Pa were used for selected experiments by adjusting the effective pumping speed of

the system. Temperature measurements were performed with a dual channel infrared

23


Quartz cupAntenna

Permanent magnet

ECR regions

Electric field

Magnetic field

Figure 3.3: Sketch of the ECR plasma source.

pyrometer, which was calibrated against a thermocouple assuming the emissivity ε as

only free parameter [76]. The experiments were carried out by applying negative high

voltage pulses of -10 kV and a pulse length of 15 µs to the substrate holder for process

times between 15 min and 2 h. The process temperature was varied between 230 and

580 ◦C by changing the pulse frequency from 0.7 to 3.0 kHz. In addition, in selected

experiments, an external infrared heating system was used. This enables to maintain

the substrate temperature without increasing the ion bombardment and the energy

deposition during the PIII pulses [76].

3.3 Analysis

3.3.1 Secondary Ion Mass Spectroscopy (SIMS)

Secondary ion mass spectroscopy (SIMS) is an analytical technique to study the com-

position of thin films [77]. A primary ion beam with energies between 0.5 and 20 keV

is focused on the surface and generates secondary particles by sputtering. As the main

part of these sputtered particles is neutral only a small fraction of ionized atoms (10−1 -

10−5) can be mass analyzed to provide information about the elemental composition

of the surface. Using reference standards it is possible to quantify the atomic concen-

tration.

In general, three different types of mass analyzers are used to separate the secondary

24

3.3. ANALYSIS

ions according to their mass-to-charge ratio: quadrupole analyzers, magnetic sector an-

alyzers and time-of-flight analyzers as in the case of this work. Afterwards, the ions are

counted with a detector. Time-of-flight secondary ion mass spectroscopy (TOF-SIMS)

usually uses a pulsed ion beam with an energy between 10 to 25 keV which is focused

on the sample surface. The sputtered ions are extracted and accelerated towards the

detector by an electric field. By measuring their mass-to-charge ratio and their time of

flight between the sample and detector, it is possible to record a mass spectrum and

determine the elemental composition of the sample.

Whereas static SIMS allows elemental analysis of atomic monolayers on the surface,

dynamic SIMS gives information about the depth distribution of several elements in

the bulk. By eroding the surface with a second ion beam and recording sequential

SIMS spectra the intensity of a given mass signal as a function of time can be ob-

tained. Assuming a linear sputter rate the intensity-time profile can be converted into

an intensity-depth profile providing that the depth of the resulting sputter crater is

known.

SIMS is a very sensitive method and it is able to analyze most of the elements with

ppm sensitivities. The lateral resolution is typically between 0.1 -1 µm whereas depth

resolutions during dynamic SIMS are generally 1 - 10 nm. The main problem existent

in SIMS measurements is the so-called ”matrix-effect”: the proportion of ions in the

sputtered flux varies considerably depending on the specific element/matrix combina-

tion studied. Moreover, a loss in depth resolution may be evidenced due to an increase

in surface roughness because of sputtering.

SIMS measurements in this work were performed using a 15 keV Ga+ beam and a

2 keV O+2 beam for sputtering, respectively. The sputter rate was determined from the

crater depth as obtained from profilometer measurements.

3.3.2 Glow Discharge Optical Emission Spectroscopy (GDOS)

In order to convert the SIMS signal intensities into atomic concentrations glow discharge

optical emission spectroscopy (GDOS) [78] was used as reference method. The principle

of GDOS is the analysis of light emitted by excited atoms after sputtering from a sample

due to glow discharge. The basic device consists of a glow discharge cell which is floated

with argon gas at a pressure between 100 and 1000 Pa. By applying a dc voltage of

about 1000 V between anode and the sample (cathode) a glow discharge is established

and material is removed from the sample by argon sputtering. Thus, the sputtered

neutral atoms are excited in the plasma and their characteristic light is collected by

25


an optical spectrometer. The intensity of the light is related to the number of emitted

atoms whereas its wavelength is characteristic of their nature.

Due to its high sputtering rate (10 - 100 nm/s [79]) GDOS is an ideal method for depth

profiling between 1 - 15 µm and it is used commercially for surface and coating analysis.

In contrast to SIMS, matrix effects only play a minor role and less non-linearities during

depth to time conversion exist. However, the depth resolution in comparison to SIMS

is poor. Another weakness of GDOS is the lack of spatial resolution as the analysis

area is millimeters in diameter.

For SIMS calibration selected samples were measured with GDOS with a voltage of

900 V and a current of 18 mA. The sputter area on the sample was restricted by the

tubular anode to a circular surface of 2.5 mm in diameter.

3.3.3 Inductively Coupled Plasma Optical Emission Spectroscopy

(ICP-OES)

A powerful method for detecting traces of metals in samples is inductively coupled

plasma optical emission spectroscopy (ICP-OES) [80, 81]. It is based on excitation of

sample atoms with plasma followed by optical spectroscopy. The principle is similar to

GDOS with the difference that liquid samples instead of solid samples are investigated.

ICP-OES uses a coupling coil to transmit radio frequency to argon gas, producing

an argon plasma. The sample which must usually be in liquid form is sprayed as an

aerosol into the plasma. The plasma removes any remaining solvent and causes sample

atomization followed by ionization. The excited atoms and ions emit electromagnetic

radiation with characteristic wavelengths of the respective elements which is analyzed

in an optical spectrometer. Hence, conclusions on the existence of elements and their

concentrations within the sample can be drawn.

ICP-OES enables one to analyze almost all the elements in the periodic table. It has

a wide dynamic concentration range and can measure elements from traces up to high

concentrations. Detection limits for most elements are in the range of micrograms per

liter to milligrams per liter. In addition, ICP-OES is a very fast technique since mul-

tielemental quantitative analysis can be carried out.

ICP-OES was performed with a plasma source operating at 1400 W excited by ra-

dio frequency of 27.12 MHz. The wavelengths corresponding to cobalt (228.616 nm),

chromium (267.716 nm) and nickel (231.604 nm) were investigated.

26

3.3. ANALYSIS

3.3.4 X-Ray Diffraction Analysis (XRD)

X-ray diffraction analysis (XRD) is a non-destructive technique for the identification

of atomic structures of non-amorphous materials [82, 83].

X-rays are generated in an X-ray source consisting of a cathode ray tube with a heated

filament which produces electrons. These electrons are accelerated towards a metal tar-

get (commonly made of copper, cobalt or molybdenum) where they create X-rays with

characteristic wavelengths. These X-rays are filtered to a single wavelength, collimated

and directed to the sample. On the sample surface the incident X-rays are diffracted

by suitably oriented crystallographic planes towards the detector which records and

processes the signal.

Geometrical considerations show that the scattering angles corresponding to diffracted

intensity maxima can be described by Bragg’s equation

nλ = 2dhkl sin θ (n = 1, 2, 3, ...), (3.1)

with the interplanar spacing dhkl from a set of hkl planes, the wavelength of the X-rays

λ, the angle of diffraction θ and the diffraction order n.

Depending on the aim of the investigation and the properties of the sample, differ-

ent experimental methods exist. For polycrystalline and powder samples the basic

measurement geometry is the θ-2θ geometry where θ is the incident angle and 2θ the

angle of the detector with respect to the incident beam. Additionally, in the so-called

222

Detector

X-ray source

Diffractometer circle

Focusing circle

Sample

Figure 3.4: Geometry of an X-ray diffractometer in Bragg-Brentano geometry.

Bragg-Brentano mode the X-ray source and the detector are always situated on the

27


same circle (see figure 3.4). The sample surface is tangential to the focusing circle and

the X-ray source and the detector are positioned on the diffractometer circle where it

intersects the focusing circle.

The diffraction pattern is obtained by varying the incidence angle of the incoming

X-ray beam and the scattering angle simultaneously while measuring the scattered in-

tensity as a function of the latter. In Bragg-Brentano geometry all reflections occur

simultaneously due to random orientation of the crystallites. But in contrast to other

geometries (like Debye-Scherrer or Seemann-Bohlin [84, 85]), only one reflection can be

well focused. This focusing condition is fulfilled when the angles between the sample

surface and incident and diffracted rays are equal.

Information about the crystal lattice, stress as well as the size of the crystals is ob-

tained from the characteristic peak positions and their half width. Depending on the

size of the angle of the incident beam penetration depths up to a few micrometers are

possible.

XRD patterns presented in this work were performed using CuKα radiation with the

wavelength 1.54056 A in standard θ-2θ geometry.

3.3.5 Scanning Electron Microscopy (SEM) and Energy Dispersive

X-Ray Analysis (EDX)

One method to investigate and to visualize surface structures is scanning electron mi-

croscopy (SEM) [86]. A finely focused electron beam with an energy up to 30 keV scans

across the sample surface and generates secondary electrons, backscattered electrons,

and characteristic X-rays. These signals are collected by different detectors to form

images of the sample.

The most common imaging mode uses secondary electrons which are produced by in-

elastic scattering with beam electrons in the surface near regions down to about 50 nm.

Since their yield depends strongly on the local surface inclination an excellent topogra-

phy contrast can be achieved. Another mode, backscattered electron imaging, is based

on elastic scattering interactions down to depth of 100 nm. They give information

about the elemental distribution as the probability of backscattering depends on the

atomic mass of the sample elements. Finally, the electron beam generates X-rays from

the sample’s atoms within a depth of one micrometer as well. The energy of each X-ray

photon is characteristic of the elements which produced it. Thus, the so-called energy

dispersive X-ray analysis (EDX) provides a further method for elemental identification.

The resolution of SEM is essentially determined by the electron beam spot size which

28

3.3. ANALYSIS

can be adjusted to values between 5 and 20 nm. Hence, a resolution limit down to a

few nanometers can be achieved.

SEM micrographs used in this work were taken by a scanning electron microscope

equipped with a Schottky field emission cathode, possible beam voltages between 0.5

and 50 kV and a resolution limit of approximately 3 nm.

3.3.6 Hardness and Wear Measurements

Indentation hardness tests and wear measurements are a useful tool to characterize the

mechanical properties of a material [87–89].

Nano and micro indentation play an important role for testing surfaces and thin films

in micro- and nanometer ranges. A tip with well-defined shape and mechanical prop-

erties is pressed into the sample with increasing load. The resulting hardness can be

calculated using the contact area between indenter and sample, the applied load as well

as the depth of penetration. Unlike traditional static indentation, in dynamic indenta-

tion the contact area is determined during the indentation phase instead of afterwards.

The recorded load-displacement curves give information about the mechanical proper-

ties such as hardness and elastic modulus. The advantage of this method is that the

penetration depth is easier to determine and external influences are at least partially

excluded.

In order to study the wear behavior a wear test has to be employed. The most common

wear test setup is the so-called pin-on-disc configuration in which a pin with a fixed

load oscillates over a flat disc. By measuring the penetration depth before and after

the wear test (e.g. with a profilometer) the wear rate can be determined.

The hardness measurements presented in this work were performed using a nanoin-

dentation setup with a standard three-sided pyramidal indentation body (Berkovich

indenter) and applied loads ranging from 3 to 20 mN. An oscillating pin-on-disc con-

figuration with an alumina ball with a diameter of 4.8 mm was used for the wear tests.

The applied load was 1 N.

3.3.7 Electrochemical Impedance Spectroscopy (EIS)

Electrochemical impedance spectroscopy (EIS) including potentiodynamic polarization

scans allow considerable information on electrochemical processes such as corrosion be-

havior and passivity [90]. The experimental setup comprises a corrosion cell with the

29


specimen as electrode. Applying a potential a current is caused to flow through the

cell due to the movements of ions through the electrolyte. The current indicates which

electrochemical processes actually take place at the anode and cathode, as well as their

rate.

Potentiodynamic polarization scans are based on slowly increasing the applied poten-

tial while measuring the resulting current. The dependency of the current density on

the potential describes the passivity behavior of the investigated specimen. In con-

trast, EIS is characterized by applying a sinusoidal potential (up to 10 mV) with a

certain frequency range and measuring the resulting current. The relationship between

potential and system response gives information about the complex impedance of the

system. Varying the excitation frequency over a broad frequency range from 1 µHz to

1 MHz an impedance spectrum can be taken. These diagrams give evidence about the

basic electrochemical processes of the corroding system and conclusions about charge

transfer, passivity and corrosion rate can be drawn. The main advantage of EIS is that

only small perturbation signals (small potentials) which do not disturb the electrode

properties need to be used.

Electrochemical experiments presented in this work were performed in Ringer solution

[91], saturated with atmospheric oxygen. The corrosion cell with a three electrode

set-up consisted of an Ag/AgCl reference, a Pt counter electrode and the specimen as

working electrode. The electrolyte temperature was maintained at 37.2 ± 0.2 ◦C.

30

4 Characterization of Nitrogen

Diffusion

The following chapter deals with nitrogen diffusion in CoCr alloys and austenitic stain-

less steel and the corresponding evolution of the nitrided layers. Some general aspects

about the method of depth profiling are presented, followed by results concerning the

dependency of the layer thickness on process time and temperature.

The presented experiments in this section were performed with a fixed pulse voltage of

-10 kV and a pulse length of 15 µs in a temperature range between 230 and 580 ◦C and

process times of 2 h. The temperature was adjusted by equilibrium frequencies between

0.5 and 4.5 kHz. Additional experiments with an external infrared heating system were

carried out to analyze the influence of the surface roughness on depth profiling: The

temperature was maintained at 420 ◦C for 2 h while the pulse frequency was varied

between 0 to 1.25 kHz resulting in an ion bombardment of up to 4×1018 at/cm2. Prior

to the implantations, nitrogen sputter cleaning was performed for 5 min at -10 kV and

1.5 kHz to remove the native oxide layer. (Experiments show that nitrogen cleaning

at these process conditions does not lead to a noticeable layer in the samples.) As the

present experimental equipment offers an easier handling with nitrogen gas, nitrogen

was chosen as sputtering gas instead of conventionally used argon.

The temporal evolution of the layer thickness was investigated at different tempera-

tures:

(i) at low temperature (450 and 320 ◦C) using initial heating frequencies of 4 and 5 kHz

for 15 min up to 450 ◦C (CoCr) and 320 ◦C (stainless steel), respectively, followed by

PIII at equilibrium frequency of 0.7 and 1.3 kHz.

(ii) The temporal development at higher temperatures (580 ◦C) was achieved with 1 h

preheating (external heating), 10 min sputter cleaning and subsequent PIII at 2.5 kHz,

-10 kV and process times between 5 and 115 min.

The samples were characterized with SIMS, GDOS, and laser profilometry. Scanning

electron microscope (SEM) investigations of microstructural cross-sections of implanted

31

4. CHARACTERIZATION OF NITROGEN DIFFUSION

samples were metallographically prepared and physically etched with a 1 keV Ar+ ion

beam.

It has to be mentioned that the investigated CoCr alloys and stainless steels (see table

3.1 in section 3.1) show similar behavior within each alloy class. Therefore, for the sake

of clarity only one or two selected alloys are presented in detail which exemplify the

process characteristics for the related alloy class.

4.1 Depth Profiling

To investigate the diffusion behavior nitrogen depth profiling with secondary ions mass

spectroscopy was performed. This method represents a central analyzing method in

0 1 2 3 4100

101

102

103

104

105

106

Fe

Co

CrN

Mn

Cr

NO

O

C

N

Cou

nt R

ate

Depth [µm]

Figure 4.1: SIMS profiles ofalloy HS188, implanted with10 kV at 400 ◦C.

this work. Therefore, its results and several details concerning the calibration with

glow discharge optical emission spectroscopy are presented. Figure 4.1 depicts typical

SIMS spectra obtained from a CoCr sample (alloy HS188) implanted at a temperature

of 400 ◦C. As can be seen, the count rates of different elements vary by more than 5

orders of magnitude, and differences in ionization efficiencies can be subsumed from the

data. The nominal Co/Cr ratio of 1.6:1 in the sample is converted to a Co+/Cr+ signal

ratio of 1:5. Especially N+ and NO+ ions are suppressed, whereas CrN+ shows a large

background due to mass interferences. For quantification of such elemental profiles,

additional calibrations have to be performed. Figure 4.2 shows a depth profile of the

32

4.1. DEPTH PROFILING

same CoCr alloy after PIII at 445 ◦C, as obtained from GDOS, together with the cor-

responding profile measured by SIMS. Except for the region directly below the surface,

Figure 4.2: Comparison ofnitrogen profiles after PIII at445 ◦C of alloy HS188 mea-sured by GDOS and SIMS.

0 1 2 3 40

5

10

15

20

25

30

35

40

0.0

1.0x10-4

2.0x10-4

3.0x10-4

4.0x10-4

5.0x10-4

Nitr

ogen

Con

cent

ratio

n [a

t.%]

Depth [µm]

GDOS SIMS

N/C

r Rat

io

very good agreement between GDOS and SIMS was obtained. Hence, the N+/Cr+

ratio from SIMS applied to the absolute calibration with GDOS from figure 4.2, was

used to define the nitrogen content in the samples throughout this work. Additionally,

matrix effects were found for the ionization efficiencies of the Cr+ and Mn+, and their

intensities decline with the N+ signal, whereas Co+ and Fe+ exhibit nearly constant

signal intensity, independent of the nitrogen count rate. Comparing different GDOS

profiles (not shown) the surface concentration reaches values between 30 and 35 at.%,

nearly independent of the process temperature. The same calibration was carried out

with stainless steel where a slightly higher surface concentration around 40 at.% was

obtained.

Depth calibration was obtained by linear conversion of sputter time into depth-scale

using the total crater depth. As checked by variations of the sputter time no differ-

ences of the sputter rate in the nitrided alloy compared to the base alloy were found in

contrast to other material systems like oxygen in NiTi with independent sputter rates

for the oxide layer and the bulk material [92].

Selected typical SIMS profiles for stainless steel and CoCr alloys are presented and

compared in figure 4.3 a) and b). For both classes of alloys thermally activated diffu-

sion is observed. But, as can be seen, a much faster diffusion of nitrogen in stainless

steel is obtained. Whereas the diffusion depth reaches values up to 4.5 µm in stainless

steel (alloy 304) at 400 ◦C the diffusion depth in CoCr alloys (alloy HS188) is about

2.5 µm for the same temperature. Nevertheless, the nitrogen depth profiles deviate

from the classical erfc-shape showing a nearly rectangular shape with an initially fast

33


diffusion in the implanted layer and a following slow diffusion in the bulk material.

This behavior is well-known for austenitic stainless steel [50]. A model assuming trap-

ping of nitrogen at chromium sites was developed to explain this behavior [62, 93–95].

Once all chromium sites are occupied, the additional incoming nitrogen can diffuse

quickly through the saturated highly nitrogen enriched layer. Combining this model

with the assumption of a concentration dependent diffusion coefficient [50, 96, 97], a

general model describing the nitrogen diffusion can be established [98]. However, one

chromium atom is supposed to trap more than one nitrogen atom as the chemical com-

positions (see table 3.1) show a chromium content of 15 - 25 at.% at high nitrogen

concentrations of 35 - 40 at.%. Therefore, the above mentioned model is not satisfying

and extensions have to be done (see chapter 6.1).

A nitrogen concentration of 2 at.% was used to define the layer thickness as well as the

resulting diffusion coefficients throughout this work. This layer thickness corresponds

0.0 0.5 1.0 1.5 2.0 2.50

10

20

30

40

50

60

70

230 °C 300 °C 340 °C 370 °C 400 °C

Depth [µm]

Nitr

ogen

[at.%

]

a) CoCr (HS188)

0 1 2 3 40

10

20

30

40

50

60

70b) Stainless Steel (304)

Nitr

ogen

[at.%

]

Depth [µm]

Figure 4.3: Typical SIMS depth profiles of a) CoCr (alloy HS188) and b) stainless steel(alloy 304) implanted with PIII at different temperatures for 2 h.

to the transition between the steep decline in the nitrogen concentration and the long

tail indicative of the native nitrogen inside the samples. For comparison, an SEM cross-

section of nitrided CoCr is shown in figure 4.4. This sample corresponds to the SIMS

depth profile implanted at 370 ◦C in figure 4.3 a). As can be seen the layer thickness of

about 1.3 µm obtained from SIMS agrees very well with the measured layer thickness

of the cross-section. The sharp transition from nitrided to bulk material observable

in the cross-section is in contrast to the transition zone with the declining nitrogen

concentration of the SIMS depth profiles. Up to now this deconvolution of the profiles

has not been mentioned in literature yet.

34

4.1. DEPTH PROFILING

Figure 4.4: Cross-section of CoCr al-loy HS188 implanted at 370 ◦C for 2 h.The nitrided layer is indicated with anarrow.

2 µm2 µm2 µm

Figure 4.5: Comparison ofroughness evolution as func-tion of fluence during PIII ni-triding of stainless steel (304)and CoCr (HS188) at 420 ◦Cfor 2 h and subsequent SIMSdepth profiling. The finalcrater depth in each casewas about 2 µm. The oxy-gen ion fluence of SIMS was3 × 1018 at/cm-2.

0 1x1018 2x1018 3x1018 4x10180

200

400

600

800

1000Steel CoCr

surfaceSIMS crater

Rou

ghne

ss R

q [n

m]

Fluence [at/cm2]

Furthermore, neither stainless steel nor CoCr show grain oriented diffusion in the SEM

micrographs of the cross-sections as it can be found in literature for Ni and some Fe

base alloys [99, 100].

A parameter influencing the shape of the depth profiles is the surface roughness. During

PIII processing an increase of the surface roughness Rq (root-mean-squared roughness)

is encountered, growing linearly with incident ion fluence at constant temperature,

voltage and process time from about 10 nm (similar to the roughness after polishing)

to close to 200 nm (see figure 4.5). When comparing the surface roughness after PIII

with the roughness in the crater of the SIMS measurement, an additive behavior can be

interfered from the data. The outlying data point for stainless steel at 1.5×1018 at/cm2

in figure 4.5 is most likely the result of the finite crater size of 300 × 300 µm2 and an

unusual grain distribution inside the crater. In addition, the grain structure remains

unmodified. Therefore, the roughening due to nitrogen insertion on interstitial sites

35


seems to be independent of the particular matrix (similar to the absolute sputter yield).

As a result of the above investigations the sharpness of the SIMS profiles will be af-

fected. This means that the depth resolution degrades after PIII at high fluences

which will be amplified during depth profiling. Hence, this additional broadening of

the leading edges of the profiles has to be considered when developing diffusion models

to explain the anomalous nitrogen diffusivity.

4.2 Temperature and Temporal Dependency of the

Layer Thickness

In the following, a layer growth model is assumed for analysis of the diffusion process,

instead of pure diffusional nitrogen redistribution. The diffusion coefficients derived

from the depth profiles considering a thermally activated diffusion with the function-

ality D = x2/4t (x - diffusion length) are plotted against the reciprocal temperature

in figure 4.6 in an Arrhenius plot. It has to be mentioned that at first the diffusivity

in the temperature range up to about 450 ◦C is investigated. The diffusion data of all

1.4 1.5 1.6 1.7 1.8 1.9 2.010-14

10-13

10-12

10-11

10-10

450 400 350 300 250

Steel304316Ti

CoCrL605HS188SY21

Diff

usio

n C

oeffi

cien

t [cm

2 /s]

1/Temperature [1000/K]

Temperature [°C]

Figure 4.6: Arrhenius plotof resulting diffusion coeffi-cients as a function of tem-perature for the investigatedalloys. The solid lines cor-respond to calculated activa-tion energies fitted to the datapoints.

alloys of stainless steel and CoCr show a quite similar temperature dependence, but

with absolute values for stainless steel being larger by an order of magnitude than for

CoCr. Fitting the data points the activation energies can be calculated to be between

36

4.2. TEMPERATURE AND TEMPORAL DEPENDENCY

0.9 - 1.1 eV in this temperature range. Therefore, since the obtained activation ener-

gies are relatively low it is assumed that the nitrogen diffuses towards the bulk along

interstitial pathways, leaving the original grain structure unmodified and finally occu-

pying primarily interstitial sites. Additionally, three-dimensional SIMS analysis shows

that diffusion along grain boundaries is not significantly faster than diffusion across

the grains.

The layer thickness of two CoCr alloys (L605 and HS188) implanted at 450 ◦C is plot-

ted in figure 4.7 for different process times. The experiments were performed using

Figure 4.7: Development ofthe layer thickness of CoCr(alloy L605 and HS188)at 450 ◦C as a function ofprocess time. The conti-nous lines represent a fitaccording to the experimentaldata using equation (4.1) withD = (9.5 ± 0.2) × 10-12 cm2/sand vs = 0.13 nm/s. Thedashed lines neglect surfacesputtering.

0 20 40 60 80 100 120 1400

1

2

3

4

5

6

Time [min]

L605 HS188Experiment v

s=0.13 nm/s

vs=0

Laye

r Thi

ckne

ss [µ

m]

fast initial heating to achieve the desired temperatures. Therefore, the time necessary

to heat the sample to this temperature, derived from temperature measurements and

checked with heat balance calculation [76] was subtracted from the total process time.

As can be seen, the evolution of the layer thickness x is in very good agreement with an

inverse parabolic growth law x ∝ t1/2 for both types of alloys. The diffusion coefficient,

determined from this law, of D = (9.5± 0.2)× 10−12 cm2/s is nearly identical for the

alloys L605 and HS188.

Additionally, the impinging ions lead to a sputtering of the surface and also influence

the nitriding kinetics. Whereas the diffusion depth increases with the square root of

time, surface sputtering scales linearly with time. Hence, an additional modification by

inserting a linear term describing this sputtering has to be performed and the growth

of the layer thickness x results in

dx

dt=

2D

x− vs (4.1)

37


where vs describes the surface recession velocity [101, 102]. Fitting the data points

in figure 4.7 according to equation (4.1) (continous lines), a sputter velocity of about

0.13 nm/s is obtained which is in good agreement with sputter yields from TRIM cal-

culations. In contrast, regarding diffusion alone, slightly higher layer thicknesses than

actually measured are achieved (dashed lines). However, as can be seen, the amount

of sputtered material is relatively small for the used process times, and it contributes

noticeably only at process times longer than 2 h. Comparing the surface recession vs

of CoCr with respective values of stainless steel (at 320 ◦C) similar results are found.

Therefore, no correction of diffusivity data obtained at constant and short process times

is necessary.

However, the data presented above is limited to moderate temperatures. Therefore, the

temporal evolution of the layer thickness at low temperatures (320 and 450 ◦C, respec-

tively) is compared to that at high temperatures (580 ◦C) and presented in figure 4.8

for a) CoCr and b) stainless steel. As already shown in figure 4.7 the layer growth at

0 1 2 3 4 5 6 7 8 9 10 110

1

2

3

4

5

6 High temperature (580 °C) Low temperature (450 °C)

Laye

r Thi

ckne

ss [µ

m]

Time1/2 [min1/2]

a) CoCr (HS188)

0 1 2 3 4 5 6 7 8 9 10 11 120

5

10

15

20

25

30

35 High temperature (580 °C) Low temperature (320 °C)

Laye

r Thi

ckne

ss [µ

m]

Time1/2 [min1/2]

b) Stainless Steel (316Ti)

Figure 4.8: Temporal evolution of layer thickness at low and high process temperatures ofa) CoCr (HS188) and b) stainless steel (316Ti).

low temperatures agrees with an inverse parabolic growth law (Here, sputtering is ne-

glected as process time does not exceed 2 h.). However, at high temperatures (580 ◦C)

CoCr exhibits a reduced diffusivity for long process times (≥ 1 h) deviating from the

expected behavior whereas stainless steel shows an increased diffusivity up to a factor

of 1.5. The reduction of the layer thickness in CoCr cannot be explained by sputtering

of the surface since the effect of sputtering is quite small and will result in a decrease

of the layer thickness of only about 120 nm, assuming a sputter recession velocity of

0.13 nm/s as calculated above. Here, a different explanation must be found to explain

38

4.3. SUMMARY OF RESULTS

the deviations at high nitriding temperatures in CoCr and stainless steel, with either

the diffusion model not applicable or the occurrence of a phase transition as the most

plausible candidates. In section 6.2 the phase composition will be studied at different

temperatures and process times in order to derive a model to explain the changes in

diffusivity at elevated temperatures.

4.3 Summary of Results

In general, CoCr and stainless steel show similar nitriding behavior within each alloy

class.

However, some differences do exist as well. The main results concerning nitrogen diffu-

sion in CoCr and stainless steel after nitrogen insertion with PIII can be summarized

as follows:

• Calibration of SIMS depth profiles with GDOS is a reasonable method for the

quantification of concentrations as the measured profiles agree well. In addition,

the layer thickness obtained by SEM cross-sections is in good accordance with

a nitrogen concentration of about 2 at.% used for determination of the layer

thickness in the SIMS profiles.

• Both alloy types show similar profile shapes deviating from the classical error-

function shape and indicating similar diffusion mechanisms: The plateau-type

profiles slowly decrease from the surface followed by a rather sharp leading edge at

the end of the profiles. Hence, diffusion models, taking concentration dependent

diffusion and trapping mechanisms into account that normally describe nitriding

of stainless steel, can be applied for CoCr as well. However, extensions have to

be done as the chromium content does not correspond to the maximum nitrogen

concentration in the alloys.

• A significant influence of the roughness on the sharpness of the depth profiles

is observed, depending on the initial roughness after surface modification before

analysis and the total sputter ion fluence during SIMS analysis. Therefore, the

depth resolution decreases with increasing layer thickness since high ion fluences

are necessary for SIMS profiling.

• The layer formation is thermally activated. The diffusivity of nitrogen is one

order of magnitude higher in stainless steel in comparison to CoCr leading to

39


layer thicknesses of up to several micrometers within some hours of process time.

Activation energies in the range of 0.9 - 1.1 eV for process temperatures up

to 450 ◦C are found indicating interstitial nitrogen diffusion. Constant surface

concentrations of nitrogen between 30 and 40 at.% are observed.

• At low process temperatures (≤ 450 ◦C), the layer thickness scales with the square

root of time for CoCr and stainless steel.

However, at process times longer than 2 h the influence of surface erosion increases

leading to a reduced effective layer thickness than would be expected without

sputtering.

• At process temperatures above 450 ◦C and process times more than 2 h, deviations

from the parabolic growth law are observed: Whereas CoCr exhibits a reduced

diffusivity of a factor of 0.6 (process time 2 h) an increase in diffusivity in stainless

steel of a factor of 1.5 is observed.

40

5 Influence of Gas and Plasma on

Adsorption, Incorporation and

Desorption Processes

The following chapter is dedicated to gas- and plasma-surface interactions. The main

aim is to understand adsorption, desorption and incorporation processes during PIII

which decisively influence the resulting nitrogen layer thicknesses. In the first part the

amount of nitrogen in the layer is compared with the incident ion dose whereas in the

second part the effect of the background pressure on the layer thickness is investigated.

In the third part the influences of nitrogen gas molecules, plasma and energetic ions

are separated and analyzed independently from each other. At the end the role of

the native surface oxide layer is studied and compared with oxide layers produced by

oxygen PIII.

The presented experiments were performed with nitrogen PIII at a pulse voltage of

-10 kV for 2 h and temperatures between 230 and 435 ◦C by changing the pulse

frequency. As in all PIII experiments the plasma was generated with one plasma

source operating at 150 W (plasma density ne = 1.6× 1010cm−3, electron temperature

Te =1.3 eV). At a nitrogen gas flow rate of 150 sccm the resulting pressure during the

experiments was 0.53 Pa. Additionally, process pressures of 0.33 and 0.82 Pa were used

for selected experiments by adjusting the effective pumping speed of the system.

To separate the influence of gas, plasma and energetic ions, further plasma and gas ni-

triding experiments were performed in the same chamber. Thus, for plasma nitriding,

the plasma was generated with three ECR plasma sources at a total power of 450 W and

a gas flow of 200 sccm which result in a working pressure of 0.7 Pa (ne = 4.4×1010cm−3,

Te =1.3 eV). The pressure and self bias are, however, much lower than in conventional

plasma nitriding processes. Additionally, the samples were externally heated for three

hours in vacuum - one hour without plasma to achieve the desired temperature and

two hours with plasma. No discernible temperature increase of the substrate was ob-

41

5. INFLUENCE OF GAS AND PLASMA

served after igniting the plasma. The gas nitriding experiments were performed with

the same gas flow, process time, heat treatment and working pressure as the plasma

nitriding experiments. However, the samples were exposed only to nitrogen gas. The

PIII sample series which is compared with the plasma and gas nitriding samples was

also preheated with the external heating in vacuum for 1 hour. In this work the terms

“plasma nitriding” and “gas nitriding” are used for the above mentioned treatments in

contrast to commercial gas and plasma nitriding methods presented in section 3.1.

In addition, to study the influence of oxide surface layers, some of the samples were

treated with oxygen PIII at a voltage of -12.5 kV and a temperature of 520 ◦C for

1 hour. Finally, sputtering cleaning experiments were performed with argon and nitro-

gen for 5 min, -10 kV and 1.5 kHz and 2.0 kHz, respectively.

5.1 Retained and Incident Fluence

To get an overview about the amount of nitrogen incorporated in the resulting layer the

retained and incident fluences were calculated. The incident fluences were derived from

heat balance calculations neglecting any heat conduction along the substrate holder as

radiative cooling dominates in the investigated temperature range. Thus, the local heat

balance can be approximated by the following equation [103]:

dTsdtcs = Q = SfUCpulse − 2σ(εsT

4s − ε0T 4

0 ), (5.1)

with the temperature Ts and heat capacity cs of the substrate holder, the pulse rep-

etition rate f , the pulse voltage U , the Stefan-Boltzmann constant σ, the emissivity

of the substrate and the environment ε and the temperature of the environment T0.

S = Sn/(Sn + Se) is a correction factor which considers that only a fraction of the ion

energy is transformed into thermal energy (Sn and Se are the nuclear and electronic

energy loss rates [32]). The first term on the right side of equation (5.1) describes the

incoming ion energy flux density with the charge density Cpulse integrated over one

high voltage pulse (see equation (2.6))

Cpulse =

∫Jidt =

∫0.6 en0

(dxsdt

+ uB

)dt, (5.2)

whereas the second term stands for radiative cooling. At equilibrium the change of

heat Q is zero as the ion energy flux density corresponds to the radiative cooling and

42

5.1. RETAINED AND INCIDENT FLUENCE

therefore, the incident ion flux can be calculated.

In contrast, the amount of retained nitrogen in the layer was derived by integrating the

nitrogen depth profiles and assuming atomic densities between (7.5 − 11) × 1022 at/cm3

corresponding to the respective densities of the alloys used as base material.

The resulting ratios of retained-to-incident fluence at different process temperatures

for CoCr and stainless steel are presented in figure 5.1. As can be seen, the ratio in-

Figure 5.1: Ratio ofretained-to-incident fluenceof CoCr (alloy HS188) andstainless steel (alloy 304)after PIII at a pulse voltageof -10 kV, different temper-atures and process times of2 h. The dotted line indicatesa ration of one, i.e. the wholeincident nitrogen is retainedin the resulting layer.

200 250 300 350 400 4500

1

2

3

4

5

6

7

8

9

10

CoCr (HS188)Steel (304)

Ret

aine

d/In

cide

nt F

luen

ce

Process Temperature [°C]

creases with increasing process temperature. Whereas at low temperatures the amount

of retained nitrogen of about 50 % is smaller than the implanted fluence, the contrary

situation is found at higher temperatures with retained amounts up to 800 % for stain-

less steel and 260 % for CoCr, respectively. Obviously, the incorporation of nitrogen is

not only influenced by the implanted ions but also by adsorption of nitrogen containing

species. At a typical repetition rate of 1 kHz, the background pressure is high enough

to allow an impinging flux of nitrogen molecules corresponding to a surface coverage of

more than one monolayer between the high voltages pulses. The duty cycle at 1 kHz

is 1.5 %, thus the assumption of this process taking place between the pulses seems

reasonable. During the voltage pulses, the energy deposition by the impinging ions

leads to a dissociation of these physisorbed molecules, with a fast chemisorption of the

resulting atoms and a subsequent diffusion inside the material. In addition, electrons

entering the depleted plasma sheath after the pulses and being accelerated to the sub-

strate may also contribute to dissociation of the molecules.

It has to be pointed out that the rate-limiting factor for the nitrogen incorporation is

the diffusion rate in conjunction with the limited surface concentration as one can see

43


when comparing figure 5.1 with the diffusivity (see figure 4.6) and the constant surface

concentration (see chapter 4.2): The low nitrogen diffusivity in CoCr leads to the maxi-

mum possible surface concentration after a short time as the transport of nitrogen from

the surface into the bulk is too slow to enable further uptake without exceeding the

maximum concentration. Additional implanted nitrogen will be removed. In contrast,

in stainless steel a much faster diffusion leads to an enhanced nitrogen transport allow-

ing a higher absolute incorporation of nitrogen from the surface. Comparing the ratio

of retained-to-incident fluence between low and high temperatures, the ratio increases

as the diffusivity augments with increasing temperature.

It has to be mentioned that determining the ion fluences with heat balance of equation

(5.1) is a first estimation only. Several further influences from the plasma like plasma

depletion due to high repetition frequencies or changing secondary emission coefficients

are excluded by this method. Sample geometry and size as well as heat conduction

effects, which should lead to minor corrections only, are neglected. Therefore, the pre-

sented results have to be considered as general tendencies, for detailed investigations

further examinations have to be done [104].

5.2 Influence of Background Pressure

To show that the background pressure is really responsible for the enhanced incorpora-

tion, experiments at different nitrogen background pressures were performed. Figure 5.2

200 250 300 350 400 4500

1

2

3

4

CoCr (HS188) 0.82 Pa 0.53 Pa 0.33 Pa

Laye

r Thi

ckne

ss [µ

m]

Process Temperature [°C]

Figure 5.2: Variation of ni-trogen background pressureresulting in different layerthicknesses of CoCr (alloyHS188).

44

5.3. COMPARISON OF GN, PN AND PIII

shows the resulting layer thicknesses of CoCr (alloy HS188) by varying the nitrogen

pressure at different temperatures. For the sake of clarity only CoCr is presented -

stainless steel shows the same behavior.

As can be seen, with increasing nitrogen pressure and process temperature, an increased

layer thickness up to a factor of two is observed. This corresponds to a larger nitrogen

content inside the samples indicating an enhancement of nitrogen incorporation. As an

increased background pressure increases the nitrogen flux towards the surface, larger

amounts of nitrogen can be incorporated. However, as mentioned in the last section,

the diffusion process is still determined by the diffusion rate in conjunction with the

limited surface concentration. Thus, the more pronounced dependence of the layer

thickness on the background pressure at higher temperatures can be attributed to the

increased diffusivity.

5.3 Comparison of Gas Nitriding (GN), Plasma

Nitriding (PN) and PIII Nitriding

As already indicated above, the incorporation of nitrogen is not only determined by

implantation of energetic ions but also by further surface processes taking place during

PIII. The aim of the following chapter is to separate these processes from each other

and to investigate them individually. As will be shown, especially dissociative nitrogen

adsorption seems to play an important role.

According to Ertl [105], dissociative nitrogen adsorption proceeds via N2 N2,ad→ 2Nad

with either molecular (N2,ad) or atomic (Nad) surface nitrogen. This process is illus-

trated by the Lennard-Jones-type potential diagram of figure 5.3. One can see that the

energy Ediss needed to dissociate a “free” N2 molecule (Ediss = 9.8 eV [106]) is much

higher in contrast to the activation energy E∗ to dissociate an adsorbed N2,ad molecule.

Depending on the surface composition and structure, different activation energies E∗

are obtained (e.g. on iron: E∗ ≈ 1 eV [105]) [107–109]. As these heterogenic catalytic

reactions play a very important role, as for example in NH3 synthesis or in cleaning

exhaust fumes, a considerable amount of literature is available, concentrating on influ-

ences of substrate type, orientation, surface steps and defects [105, 110–115].

However, here the interest lies on surface dissociation followed by incorporation into

the bulk where much less literature can be found. During PIII, several independent

species are arriving on the surface:

45


Ediss

Ead

E*

E

rN2

N2,ad

2Nad

Figure 5.3: Potential energy diagram asfunction of distance r from the surface fordissociative adsorption of nitrogen (Ediss =9.8 eV). For pure iron, the activation energyE∗ is about 1.0 eV, however, depending on thesurface structure. The adsorption energy Ead

is about 2.2 eV (after Ertl [105]).

• High energy ions with an energy of E = 10 keV corresponding to an average

nitrogen flux of JPIII = 1014 nitrogen atoms/cm2s at a process temperature of

480 ◦C (see equation (5.1)).

• Low energy ions with E < 5 eV and a flux of Jplasma = 1015 nitrogen atoms/cm2s

(see equation (2.4)). Simultaneously, electrons are bombarding the surface.

• Gas molecules with a thermal energy of E = 0.03 eV corresponding to about 50 ◦C

and a flux given by kinetic gas theory of Jgas = p/(2πmkBT ) = 1018 nitrogen

atoms/cm2s (p denotes the nitrogen partial pressure of 0.5 Pa) [116].

In the following, the influence of each species will be investigated separately by com-

paring PIII, plasma nitriding (PN) and gas nitriding (GN) and by determining their

respective contribution to the layer formation.

For PIII at 480 ◦C, figure 5.4 a) shows nitrogen depth profiles of one selected CoCr

alloy (HS188) and one selected stainless steel (316Ti) for a process time of 2 hours. As

already mentioned, a thermally activated diffusion is observed for both alloys which

is much faster in stainless steel than in CoCr. An inverse parabolic growth of the

layer thickness as a function of process time has already been established in section

4.2. When removing the high energy ion component from the process (plasma nitriding

only), differences in the qualitative behavior arise between the two alloy classes (see

figure 5.4 b). As can be seen for CoCr, a layer thickness of about 2 µm is observed at

480 ◦C, whereas no diffusion occurs in stainless steel with a nominal layer thickness of

less than 10 nm. Removing the plasma component (gas nitriding only), i.e. low energy

ions and electrons, the nitrogen incorporation and subsequent diffusion is eliminated

46

5.3. COMPARISON OF GN, PN AND PIII

0 2 4 6 8 10 120

10

20

30

40

50a) PIII Nitriding

Steel

CoCr

Nitr

ogen

[at.%

]

Depth [µm]0.00 0.01 0.5 1.0 1.5 2.0 2.50

10

20

30

40

50b) Gas and Plasma Nitriding

Steel (GN)

CoCr (GN)

Steel (PN)

CoCr (PN)

Nitr

ogen

[at.%

]

Depth [µm]

Figure 5.4: SIMS nitrogen depth profiles for stainless steel grade 316Ti and CoCr alloyHS188 after a) PIII at -10 kV and b) gas nitriding (GN) and plasma nitriding (PN), allprofiles for 2 h at 480 ◦C.

even for CoCr alloys: For pure gas nitriding at 480 ◦C, effective nitrogen containing

layers of less than 5 nm are observed for both investigated alloys.

To get an overview, figure 5.5 compares the diffusion coefficients for PIII and PN, as

calculated from the respective depth profiles for different process temperatures and dif-

ferent alloys. Diffusion coefficients of GN experiments are not shown since no diffusion

was observed. The data of PIII nitriding are taken from figure 4.6. In contrast to

Figure 5.5: Diffusion co-efficients as a function ofreciprocal temperature forthe investigated CoCr alloys(HS188, L605 and SY21)and stainless steels grades(304 and 316Ti) after PIIIand plasma nitriding. Thesolid lines are fits to the datapoints, whereas the dottedlines are indicated to guidethe eye.

1.2 1.3 1.4 1.5 1.6 1.7 1.8 1.9 2.010-17

10-16

10-15

10-14

10-13

10-12

10-11

10-10

10-9

550 500 450 400 350 300 250

Steel PIII PN304316Ti

CoCrL605HS188SY21

App

aren

t Diff

usio

n C

oeffi

cien

t [cm

2 /s]

1/Temperature [1000/K]

Temperature [°C]

PIII, very low diffusion is observed for PN of steel 316Ti with small layer thicknesses in

the order of nanometers, instead of some micrometers for PIII, and a very small tem-

47


perature dependence. For CoCr, intermediate diffusion coefficients are derived with

a strong dependence on the temperature. At the same time, a slight influence of the

alloy composition is present.

How can these effects be explained by dissociative adsorption processes? For GN, the

reaction speed is rather low [114, 117] and no nitrogen is incorporated into the material

in the present experiments. For PN at the present conditions, the species arriving at

the surface are a mixture of nitrogen molecules and low energy nitrogen molecular ions

with about 5 − 10 % atomic ions present in the plasma [118]. The energy of the

ions at the surface is about 5 eV, and therefore below the energy necessary to split

the nitrogen-nitrogen triple bond on impact as long as no interactions occur (dissocia-

tion energy of N2 is about 9.8 eV [106]). Though, as explained above, the interaction

of the nitrogen molecules with the surface reduces the energy for dissociation E∗ and

activated dissociation of the adsorbed molecules may take place. Subsequently, diffu-

sion into the bulk will occur. For pure iron, the energy barrier E∗ is about 1 eV, just

beyond the kinetic ion energy but above the energy of gas molecules. However, no in-

formation on the rate constant is available from the energy diagram. As thinner layers

than for PIII are observed, the nitrogen uptake through the surface is supposed to be

the rate-limiting factor, apparently strongly depending on the chemical composition of

the alloy. Additionally, for stainless steel, the incorporation of nitrogen is very low.

Alternatively, native oxide layers (see section 5.4) may provide an explanation for the

very small layer thicknesses after PN.

5.4 The Role of a Surface Oxide Layer

Due to their affinity to oxygen, CoCr and stainless steel are always covered with a

native oxide layer. The approximate thickness is about 2 nm at room temperature for

both types of alloys [119, 120]. Exactly this surface layer on stainless steel has been

already identified as a diffusion barrier for nitrogen [121]. For standard PIII processes,

this oxide layer will be sputtered away in less than one minute [104] leading theoreti-

cally to a small temporal delay of nitrogen diffusion and therefore, to a reduced effective

nitrogen layer thickness. Correspondingly, to observe a larger time delay, an additional

thermal oxide layer of 20 - 30 nm was produced on the surface of some samples. The

projected range of 10 keV N+2 ions is about 10 nm in stainless steel. Thus, implanta-

tion “behind” the oxide layer, directly into the bulk, can be excluded. Subsequently,

the nitriding behavior of the samples with thermal oxide was compared with that of

48

5.4. THE ROLE OF A SURFACE OXIDE LAYER

samples with only a thin native oxide on the surface.

Figure 5.6 presents such SIMS depth profiles after both, PIII and plasma nitriding.

As expected, PIII results in nitrogen enriched layers for both alloy classes, stainless

0 2 4 6 8 10 120

10

20

30

40

50a) PIII Nitriding

Steel(thermal oxide)

CoCr (thermal oxide)

Depth [µm]

Steel(native oxide) CoCr

(nat. oxide)

Nitr

ogen

[at.%

]

0.0 0.5 1.0 1.5 2.0 2.5 3.00

10

20

30

40

50b) Plasma Nitriding

CoCr (PIII oxide)

CoCr (nat. oxide)

Nitr

ogen

[at.%

]

Depth [µm]

Figure 5.6: SIMS nitrogen depth profiles for a) CoCr alloy HS188 and stainless steel grade316Ti after -10 kV PIII for 2 h at 480 ◦C and b) after plasma nitriding for 2 h at 480 ◦C ofCoCr (HS188). In each diagram the profiles are shown for samples with a native oxide layerand a thermal oxide layer. For plasma nitriding of stainless steel, again no diffusion beyond10 nm was observed.

steel and CoCr, as the surface oxides are sputtered due to the impinging ions (figure

5.6 a). For stainless steel, the time delay is clearly visible from the different layer

thickness. Using a sputter yield of 1.5 atoms/ion (obtained from TRIM calculations)

and a diffusion coefficient of 1.7× 10−11 cm2/s (compare figure 4.6) for nitrogen diffu-

sion in stainless steel, a thickness reduction of about 1.6 µm is obtained at 2 h which

is in good agreement with the observed change in the depth profile. The interesting

point is that for CoCr, the thermal oxide presents no barrier as no reduced thickness

is found. However, this thick thermal oxide of CoCr completely suppresses plasma

nitriding which is occurring for the thin native oxide, as shown in figure 5.6 b). In

the case of stainless steel, no nitrogen incorporation after plasma nitriding was found

neither with the thermal nor with the native oxide (cf. figure 5.4 b).

To elucidate differences between the thermal and the native oxide, the element enrich-

ments in the near-surface regions of both types of oxides were examined for stainless

steel and CoCr (see figure 5.7). Comparing the native and the thermal oxides, a layer-

like structure with nickel near the surface followed by chromium is visible for the native

oxides. In contrast, almost constant nickel and chromium ratios are found in the ther-

mal oxides. Thus, it is assumed that these different kinds of element enrichments in

49


0 1 2 3 4 5

0.1

1

a) Stainless Steel

native thermaloxide oxide

Cr/Fe Ni/Fe

Inte

nsity

Rat

io

Depth [nm]

0 1 2 3 4 5

1

10

native thermaloxide oxide

Cr/Co Ni/Co

b) CoCr

Inte

nsity

Rat

io

Depth [nm]

Figure 5.7: SIMS depth profiles of the near surface-regions for a) stainless steel (316Ti) andb) CoCr (HS188) with a thermal oxide layer (produced with oxygen PIII at a pulse voltage of-12.5 kV at a temperature of 520 ◦C for 1 h) and with a native oxide layer (untreated).

the surface oxides are responsible whether nitrogen is able to diffuse through the oxide

or whether it is hindered. However, further differences should be present to explain not

only the differences between the native and the thermal oxides but also the differences

between CoCr and stainless steel.

These above presented results can be explained by the following model for the nitrid-

ing process: dissociative adsorption of nitrogen molecules is active at low rates for

CoCr surfaces when enhanced by low energy ion or electron bombardment. A strong

dependency of this process on the chemical composition is present as already shown

in chapter 5.3. As a second step, transport from the surface into the material and

through an oxide layer has to occur. Again, differences in the oxide layer can influence

the transport behavior, e.g. delay or prohibit it. The distinction between a surface

process and a transport below the surface is exemplified by the different results on PIII

and plasma nitriding treatment of the CoCr alloy with a thermal oxide. Ion implanta-

tion into the thermal oxide leads to diffusion, whereas surface adsorption on a thermal

oxide does not lead to noticeable diffusion in contrast to a native oxide.

Using ion implantation, these first two steps are circumvented and the dissociated ni-

trogen molecules are directly inserted below the oxide layer (or into the surface after

sputtering the oxide) and can directly participate in the diffusion process. Thus, the

fastest diffusion and the largest layer thickness are observed here with the diffusion

as rate-limiting process and not the nitrogen incorporation and transport through the

surface. In the case of stainless steel, the barrier of a native oxide is sputtered away

50

5.4. THE ROLE OF A SURFACE OXIDE LAYER

during the first minutes, thus allowing an incorporation in the latter phase of the pro-

cess.

As a consequence of these investigations, using a preliminary sputtering step before

plasma nitriding should allow to achieve a measurable nitrogen incorporation for stain-

less steel similar to CoCr. To study the efficiency of pre-sputtering, different ion species

are compared. Figure 5.8 shows depth profiles after PN with previous sputtering using

argon and nitrogen, respectively. For comparison, CoCr is presented, too, although

Figure 5.8: Depth profiles ofCoCr (HS188) and stainlesssteel (304) after PN with pre-liminary sputtering for 5 minwith argon and nitrogen, re-spectively. PN was performedfor 2 h at 425 ◦C.

0.0 0.2 0.4 0.6 0.8 1.0 1.20

5

10

15

20

25

30

35

40

Ar N2

CoCrSteel

Nitr

ogen

[at.%

]

Depth [µm]

there is no need for sputtering cleaning as shown above. As can be seen, similar layer

thicknesses between 200 and 600 nm are obtained for each alloy class, independent of

the sputtering species.

In contrast to PIII, lower nitrogen incorporation is observed during plasma nitriding of

stainless steel than for CoCr despite removing the oxide layer. This behavior may be

explained by oxygen remnants within the residual gas. Depending on the oxygen par-

tial pressure an interplay between re-oxidation and sputtering takes place [46]. Hence,

the formation of a thin oxide layer is possible. Depending on the composition of the

alloy and the nature of the oxide, it suppresses nitrogen incorporation of stainless steel

to a greater extend than that of CoCr as already shown. Alternatively, different disso-

ciation efficiencies may be present for different chemical compositions of the alloys. In

literature it is shown that high nickel contents on the surface of stainless steel catalyzes

dissociation and prevents repassivation [122].

51



Nitriding of CoCr and stainless steel is not only influenced by implantation of ions

but also on additional processes active during PIII. The results of this chapter can be

summarized as follows:

• Plasma-surface interactions play a decisive role during ion implantation with PIII.

The following model is proposed:

Nitrogen is adsorbed in the pulse pauses and diffuses into the material. The diffu-

sion rates combined with the limited maximum surface concentration of nitrogen

in the layers are the rate-limiting factors during this process. Consequently, an

enhanced nitrogen incorporation from the plasma can be found at elevated tem-

peratures (≥ 350 ◦C) (compared to lower temperatures) and for stainless steel

(compared to CoCr) due to higher nitrogen diffusivity.

• The incorporation of nitrogen can be enhanced by increasing the nitrogen back-

ground pressure leading to an increased adsorption of nitrogen molecules on the

surface. This effect is pronounced at higher temperatures.

• Separating the influence of gas, plasma and high energetic ions (e.g. 10 keV), the

following conclusions can be drawn:

- No incorporation of nitrogen from the gas phase is found at the present

conditions as the dissociative adsorption rate is too slow.

- PN (without sputter cleaning) enables nitrogen uptake as adsorption of low

energy ions leads to activation of the surface enhancing dissociation of ad-

sorbed molecules. The following transport from the surface into the bulk

is determined by the chemical composition of the alloy and the surface ox-

ide. Thus, different nitriding behavior is found for the two alloy classes with

mediate layer thicknesses (up to 1.1 µm at 540 ◦C for 2 h) for CoCr. In

contrast, no noticeable layer is obtained for stainless steel.

- PIII circumvents these restrictions in implanting ions directly behind the

surface barrier. At the same time, the surface is simultaneously eroded by

the impinging ions. Thus, the diffusivity of nitrogen is the rate-limiting

factor of nitrogen incorporation with PIII.

To get an overview, the main processes present during GN, PN and PIII, are

summarized in figure 5.9. Besides, the respective rate-limiting factors for each

52


GN PN PIII

Adsorption

Thermaldissociation

(extremelyslow)

Activateddissociation

Transport through oxide

Implantationbehind surface

barrier(R = 10 nm)p

Sputteringof oxide

Thermally activated diffusion

Rate-limitingfactor

GN:Thermallyactivated dissociation(extremely slow)

PN:Activated dissociationand transport through oxide(composition of alloy andoxide can prohibit transport)

PIII:Diffusion rate in bulk material/nitrided layer and constantnitrogen surface concentration(oxide can only delay diffusion)

Steel

Thermaloxide

Nativeoxide

Thermaloxide

Nativeoxide

Thermaloxide

CoCr

barrier barrier

barrier barrier

barrier

Nativeoxide

barrier

barrierdiffusion

diffusion

diffusion diffusion

delayeddiffusion

Bulk

Oxide

Flux of moleculesJ = 10 at/cm s

E(N ) = 25 meVgas

2

18 2

Low energy flux ofmolecular ions

E( ) = 5 eV

J = 10 at/cm s

Nplasma

2

15 2

+

High energy fluxof molecular ions

10 keV

J = 10 at/cm s

E(N ) =PIII

2

14 2

+

Physi-sorptionduring pulsepauses anddissociationdue toimpingingions

Figure 5.9: Overview of processes during gas nitriding (GN), plasma nitriding (PN) andplasma immersion ion implantation (PIII). The rate-limiting factors for each process and therole of native and thermal oxide layers concerning diffusion in stainless steel and CoCr arelisted thereunder. The results refer to the process conditions present in this work.

process are listed.

• In addition, the thickness and nature of a surface oxide layer influence diffusivity.

The investigations with thermal and native oxide layers show the following (see

also the table of figure 5.9):

- For stainless steel, implantation into the oxide leads to a temporal delay of

nitrogen diffusion as this oxide layer has to be sputtered. In contrast, no

53


delay is found for CoCr.

- A thick thermal oxide layer prohibits nitrogen incorporation in stainless steel

and CoCr during plasma nitriding.

Therefore, different metal enrichments in the native and thermal oxides of stain-

less steel and CoCr have to play the decisive role and influence the dissociation

efficiency as well as the diffusivity.

• Sputtering cleaning enables the formation of nitrogen enriched layers during

plasma nitriding. Experiments with different ion species show that argon as

well as nitrogen are applicable for sputtering at the present conditions (sputter-

ing time of 5 min at -10 kV). Nitrogen sputtering does not lead to a noticeable

layer formation.

54

6 Phase Formation

As diffusivity of additionally inserted atoms is strongly correlated with the crystal

lattice of the host material, investigation of phase formation after nitrogen insertion is

essential to understand the diffusion mechanisms and the layer growth rates. Therefore,

in the following chapter, phase formation of CoCr and austenitic stainless steel after

nitrogen plasma immersion ion implantation will be studied. The first part presents

the formation of expanded austenite at elevated temperatures due to lattice expansion,

whereas the second part is dedicated to the decomposition of this metastable compound

at certain process conditions. Finally, in the third part, the dependency of the lattice

parameter on the nitrogen content will be studied and the validity of Vegard’s law will

be examined.

The phase characterization was performed with X-ray diffraction in θ-2θ geometry.

XRD patterns presented in section 6.1 were taken at room temperature after PIII with

a voltage of -10 kV, at temperatures between 230 and 580 ◦C and a process time of

2 h.

In contrast, to study the temporal evolution of the phase formation (section 6.2), two

sets of experiments were performed. In the first set, XRD was carried out during

annealing in a vacuum furnace at three different constant temperatures (nominal 325,

375 and 425 ◦C). The heating rate to achieve the desired temperature was 20 K/min.

Beforehand, PIII was performed at 400 ◦C for 1 h. In the second set, PIII was carried

out with different process temperatures between 15 min and 2 h after preheating up

to 580 ◦C for 1 h. After cooling down the samples to room temperatures, the X-ray

diffraction patterns were measured.

6.1 Lattice Expansion

To get an overview about phase formation after nitrogen PIII, XRD patterns of stain-

less steel and CoCr are compared in figure 6.1 for different temperatures. As can be

55

6. PHASE FORMATION

seen, the same temperature dependence is observed for both alloy classes: The base

material exhibits diffraction lines corresponding to an fcc lattice which disappear with

increasing temperature due to the increasing layer thickness of the overlying modified

layer. In contrast, the surface layer itself is characterized by an expanded fcc lattice

with clearly visible shifts of the austenite peaks towards lower diffraction angles (es-

pecially the peaks according to reflections of (111) and (200) of the fcc lattice are

pronounced as they exhibit the highest intensities). This expanded phase is found in

nitrided CoCr with similar characteristic reflections as those of the typical γN peaks

of stainless steel. With increasing temperature these peaks are shifted to lower angles

indicating a larger lattice expansion (6 - 12 %) because of the insertion of nitrogen

atoms.

However, expanded austenite is metastable [123]. When the treatment temperature

exceeds approximately 400 ◦C, the expanded austenite decomposes, the γN peaks lose

intensity and reflections corresponding to CrN and Cr2N precipitates appear. Addi-

tionally, higher diffraction intensities are found for CoCr in comparison to stainless

steel, indicating a better crystalline quality of CrxN as the chromium content is similar

in both systems. The chromium-depleted matrix resorts to a bcc (ferrite) structure for

steel which is in accordance with literature [124, 125]. In contrast, in CoCr an fcc CoNi

phase with reflections almost at the same positions as the fcc base material is found.

In an fcc crystal lattice, nitrogen atoms occupy octahedral interstitial sites [55] (see fig-

ure 6.2). These interstitial sites provide a natural limit for the total amount of nitrogen

which can be incorporated. As shown in the depth profiles (see chapter 4.1), a satu-

ration of the nitrogen surface concentration around 30 to 40 at.%, nearly independent

of the alloy or process conditions, is obtained. At the same time, no large decrease of

the concentration near the surface is observed, indicating no or only a minor activity

of the surface as a sink for nitrogen, i.e. outgasing or outdiffusion [121].

Additionally, a constant nitrogen supply rate for PIII contrasts with an inverse parabolic

layer growth: with increasing process time, an ever greater excess of nitrogen should be

accumulating in the near surface region, which is in contrast to the saturated surface

concentrations. Going back to the occupancy of possible nitrogen sites, the following

model is proposed. A ratio of 1 nitrogen atom to 2 atoms of the base alloy corresponds

to 33 at.% nitrogen. Occupying the octahedral sites in the fcc lattice, this would imply

a maximum of 2 nitrogen atoms in one (expanded) unit cell of the fcc lattice, which

appears as a reasonable assumption. Any nitrogen in excess of this limit should be in-

corporated with a much lower binding energy and should be much more easily released

into the gas phase. At the same time, the higher binding energy for the interstitially

56

6.1. LATTICE EXPANSION

0

100

200

300

35 40 45 50 60 70 80 90 100 1100

100

200

bcc

Fe

bcc

FeCr 2N

Cr 2N C

rNCrN

(111

)

bcc

FeC

rN

Cr 2N

435 °C

435 °C

555 °C

555 °C

580 °C

b) CoCr

a) Steel

(222

)(311

)

(220

)(200

)

Inte

nsity

[a.u

.]

(222

)(311

)

(220

)

CrN(2

00)

CrN(1

11)

Cr 2N

Cr 2N

Cr 2N

CrN

untreated

untreated

230 °C

230 °C

300 °C

300 °C

370 °C

370 °C

580 °C

Inte

nsity

[a.u

.]

Angle 2 [°]

Figure 6.1: XRD patterns for a) stainless steel alloy 304 and b) CoCr alloy HS188 fortemperatures between 230 and 580 ◦C. For sake of clarity, the peak positions of the expandedphase γN are not indicated.

bound nitrogen (up to 33 at.%) would reduce the probability of outdiffusion for this

class of nitrogen atoms.

Investigations on nitriding of Fe-Cr alloys [126] are along the same lines of reasoning,

establishing three different nitrogen sites in their case: precipitated as Cr-N, equilib-

57

6. PHASE FORMATION

- Fe, Co, Cr, Ni, W, ...

- N (octahedral interstitial sites)

Figure 6.2: Octahedral interstitial sites of the fcc lattice of austenitic stainless steel andCoCr.

rium saturation of the ferrite matrix and nitrogen adsorbed at grain boundaries and

nitride/matrix interfaces. In addition, when gas nitriding of thin austenitic steel foils

is performed with different nitrogen activities, a maximum occupancy of the octahe-

dral sites of 61 %, corresponding to 38 at.% nitrogen, was found [122]. By varying

the nitriding potential, different nitrogen activities could be obtained leading to dif-

ferent quantities of nitrogen in the samples. Finally, at an infinite nitrogen activity a

maximum occupancy of the octahedral sites of 61 % was obtained.

6.2 Temporal Decomposition of Expanded Austenite

6.2.1 Phase Formation after PIII during Different Annealing Times

The structural behavior of nitrided stainless steel and CoCr alloys was studied after

implantation at different annealing temperatures as function of time. Figure 6.3 shows

X-ray diffraction patterns of steel (panels a-c) and CoCr (panels d-f) during annealing

at 325, 375 and 425 ◦C for up to 8 h. Preliminarily, the samples were nitrided with PIII

at 400 ◦C for 1 h in order to form expanded austenite (γN) which can be observed in

the respective lowest patterns which were taken before annealing at room temperature

(RT).

For stainless steel at 325 ◦C (figure 6.3 a), the expanded phase decomposes after an an-

nealing time of 1.5 h and peaks corresponding to Cr2N and bcc Fe appear. During the

whole annealing process, the diffraction lines of the fcc lattice of the base material (γ)

are visible. This is in accordance with the energy of the X-rays which is high enough

to get information of the bulk material “under” the nitrided layer as layer thicknesses

between 2 - 2.5 µm were measured with secondary ion mass spectroscopy. The patterns

of steel at 375 ◦C (b) and 425 ◦C (c) show similar behavior. However, the decomposi-

58

6.2. TEMPORAL DECOMPOSITION OF EXPANDED AUSTENITE

36 38 40 42 44 46 48 50 52

400

800

1200

1600

gN

gN

Cr2N gg bcc Fe

a) Steel - 325 °C

4.5 h

1.5 h

1 h

0.5 h

0 h

RT

RT

Inte

nsi

ty[a

.u.]

Angle 2q [°]

36 38 40 42 44 46 48 50 52

200

400

600

800

gN

gN

g

bcc FegCrN/Cr2N

b) Steel - 375 °C

RT

0 h

0.5 h

1 h

1.5 h

4.5 h

RT

Inte

nsi

ty[a

.u.]

Angle 2q [°]

36 38 40 42 44 46 48 50 52

200

400

600

800

gN

gN

CrN/Cr2N

Cr2N

g

gbcc Fe

c) Steel - 425 °C

RT

0 h

0.5 h

1 h

1.5 h

4.5 h

RT

Inte

nsi

ty[a

.u.]

Angle 2q [°]

36 38 40 42 44 46 48 50 52

200

400

600

800

gN

gN

Cr2N g

g

d) CoCr - 325 °C

8 h

RT

1 h

2 h

3 h

4 h

RT

Inte

nsi

ty[a

.u.]

Angle 2q [°]

36 38 40 42 44 46 48 50 52

300

600

900

gN

gN

CrNfcc CoNig

g

Cr2NCrN/Cr

2N

e) CoCr - 375 °C

RT

6 h

1.5 h

1 h

0.5 h

0 h

RT

Inte

nsi

ty[a

.u.]

Angle 2q [°]

36 38 40 42 44 46 48 50 52

300

600

900 Cr2N

CrN/Cr2N

gN

gN

fcc CoNig

g

CrN

f) CoCr - 425 °C

RT

6 h

1.5 h

1 h

0.5 h

0 h

RT

Inte

nsi

ty[a

.u.]

Angle 2q [°]

Figure 6.3: XRD patterns of a-c) steel (alloy 304) and d-f) CoCr (alloy HS188) duringannealing at 325 ◦C, 375 ◦C and 425 ◦C up to 8 h. The samples were pretreated with PIII at400 ◦C for 1 h resulting in expanded austenite γN (see lowest pattern at room temperature).Before and after annealing a pattern was taken at room temperature (RT). The label “0 h”corresponds to the first pattern after heating to the desired temperature. γ indicates the fccphase of the base material.

59

6. PHASE FORMATION

tion of γN is much faster and the expanded phase disappears after 0.5 h (375 ◦C) and

even already during the heating phase (425 ◦C). Furthermore, the intensities of the

Cr2N peaks increase and additional diffraction lines corresponding to CrN are clearly

visible.

Annealing of nitrided CoCr leads to almost similar results (figure 6.3 d-f): The pat-

terns at room temperature before annealing show reflections corresponding to expanded

austenite which disappears with increasing annealing time. Similar to steel, the decom-

position rate is more pronounced at higher temperatures. Whereas the diffraction lines

of γN are visible up to 4 h at 325 ◦C (d), the γN phase decomposes after 0.5 h at 375 ◦C

(e) as well as during the heating phase (425 ◦C) (f). Instead of the formation of a bcc

lattice in stainless steel, in CoCr further peaks at 2θ = 44.0 ◦ and 2θ = 51.3 ◦ are

revealed corresponding to the fcc lattice of the chromium-depleted CoNi-matrix. Obvi-

ously, the decomposition of the expanded phase and the formation of CrN/Cr2N does

not only depend on temperature but also on time. In addition, whereas hcp Cr2N is

formed after a respective annealing time independent of the temperature (325 - 425 ◦C),

the formation of fcc CrN is favored at higher temperatures (375 and 425 ◦C).1 From

this time- and temperature-dependent decomposition of γN, the activation energies of

CrN and Cr2N formation could be calculated. However, a wider range of annealing

temperatures as well as an exact calibration of the actual temperature are indispens-

able to get reliable values thereof.

It has to be noted that a slight shift of the diffraction lines of about 0.3 ◦ towards lower

angles is observed during annealing for both alloy types. As can be seen, this shift

disappears when cooling down the samples to room temperature, thus indicating a lat-

tice dilation due to the temperature treatment. Additionally, in the X-ray diffraction

patterns, no peak broadening of the diffraction lines neither of Cr2N, CrN, bcc Fe nor

γ is observed as can be found by other groups [129].

Investigating the nitrided layers before and after annealing with SIMS, nearly constant

values of 2 µm for CoCr and 4 µm for stainless steel were obtained before and after

annealing of up to 8 h. This indicates that annealing at these process conditions does

not lead to any nitrogen mobility and a continuous nitrogen supply is necessary to

provide mobile nitrogen atoms.

1It is suggested that the actual annealing temperatures were about 75 ◦C higher than reported asthe precipitation of Cr2N and CrN at temperatures around 325 and 375 ◦C is unusual - even aftersome hours of annealing. In literature no precipitation of chromium nitrides in stainless steel canbe found at nitriding temperatures below ∼400 ◦C [127, 128]. Therefore, it is assumed that theactual annealing temperatures were between 400 and 500 ◦C.

60

6.2. TEMPORAL DECOMPOSITION OF EXPANDED AUSTENITE

6.2.2 Phase Formation after PIII with Different Process Times

In figure 6.4, the X-ray diffraction patterns of stainless steel (alloy 316Ti) and CoCr

(alloy HS188) after different PIII process times are presented. The corresponding layer

Figure 6.4: XRD patterns for a) stainless steel alloy 316Ti and b) CoCr alloy HS188 afterPIII at process times between 15 and 120 min at 580 ◦C. Again, for sake of clarity, the peakpositions of the expanded phase γN are not indicated.

thicknesses of theses patterns are shown in figure 4.8 in chapter 4.2. Analyzing the

XRD data, one can see that the expanded phase in stainless steel and CoCr is present

up to 30 min although the process temperature is relatively high (580 ◦C). However,

61

6. PHASE FORMATION

again for longer treatment times the γN phase decomposes, and chromium nitrides and

a chromium depleted matrix (bcc Fe and fcc CoNi, respectively) are formed. As already

seen in figure 6.3 the diffraction lines of CrN are much better visible in CoCr than in

stainless steel where Cr2N are primarily formed.

These results suggest a correlation between phase formation and nitrogen diffusivity,

where deviations from the parabolic layer growth of the layer thicknesses with increasing

process time, especially at high temperatures, were observed (figure 4.8). In stainless

steel, the increased diffusivity can be attributed to a more rapid interstitial diffusion of

nitrogen in bcc Fe than in fcc Fe [130]. Comparing the respective diffusion coefficients

of nitrogen in both phases, a good accordance with literature is found [62]. In contrast,

in CoCr, the decomposition of the expanded phase in an fcc CoNi lattice has contrary

influences on diffusivity leading to a reduced diffusivity and lower layer thicknesses

than expected. Preliminary experiments to determine the nitrogen diffusivity in pure

cobalt and nickel show that no noticeable mobility of nitrogen is present there [131].

Therefore, the absence of nitrogen mobility in CoNi is understandable.

As can be seen in the X-ray diffraction patterns after short annealing times (up to

1 h) (figure 6.3), diffraction lines of CrN/Cr2N together with the expanded phase are

present. Taking into account the temporal dependence of the decomposition of the

expanded phase, this suggests the following model. The decomposition starts from

the surface towards the bulk, as γN has been existed there for the longest time (see

figure 6.5). Therefore, for a certain period of time, two layers exist: the decomposed

expanded austenite (CrxN with either bcc Fe for steel or fcc CoNi for CoCr alloys) and

the expanded austenite thereunder which has not been decomposed yet. In the case of

steel, bcc Fe is formed with accelerated diffusion of nitrogen through this layer. Con-

sequently, the total layer thickness will increase as enough nitrogen can move deeper

and deeper into the bulk.

However, a different behavior is observed for CoCr. The formation of the fcc CoNi

phase after removing chromium and nitrogen leads to a reduced nitrogen diffusivity as

explained above. Thus, only minor amounts of nitrogen can be provided at the tran-

sition zone between nitrided layer and bulk, eventually terminating the layer growth.

Finally, with increasing time and temperature the expanded phase is completely de-

composed. This is also visible in figure 6.4 after varying PIII process times, however,

only for CoCr as the minimum layer thickness of stainless steel (∼ 4 µm after 15 min) is

too large to see any influence from layers between the decomposed expanded austenite

and the bulk.

For further quantitative evidence of this model additional experiments like in situ XRD

62

6.3. LATTICE PARAMETER AND NITROGEN CONTENT

Expanded lattice

( )gN

fcc steel bulk

( )g

Expanded lattice

( )gN

fcc CoCr bulk

( )g

Expanded lattice ( )gN

Expanded lattice ( )gN

CrN/Cr N + bcc Fe2 CrN/Cr N + fcc CoNi2

Increasing temperature and time

Stainless steel CoCr alloys

Figure 6.5: Schematic development of phase formation and diffusion of stainless steel andCoCr with increasing time and temperature. The small arrows indicate the existence andintensity of the layer growth.

during implantation are necessary. The above presented ex situ investigations only ex-

hibit some kind of “snapshots” after implantation and are only qualitative tendencies.

6.3 Lattice Parameter and Nitrogen Content

In several publications, the nitrogen concentration after nitriding of stainless steel and

other alloys is estimated by assuming a linear relation between the lattice parameter

and the fraction of occupied interstitial sites CN [99, 132–135], mostly to avoid the

measurement of the absolute nitrogen concentrations. This relation is valid for physics

in equilibrium state and it is known as Vegard’s law for interstitially dissolved atoms

([136, 137]) and can be expressed as

aγN= aγ + αCN , (6.1)

63

6. PHASE FORMATION

where aγNand aγ are the lattice parameters for the nitrogen containing and nitrogen

free γ phases, respectively, and α is the Vegard’s constant.

As nitrogen content and lattice expansion were measured separately in this work, a

comparison of the results is possible. In figure 6.6 the lattice parameters calculated

from the peak positions of the γN phases for the (111) and the (200) reflections are

plotted against the number of nitrogen atoms per 100 metal atoms. The nitrogen

concentrations were estimated from GDOS measurements from the maximum concen-

tration values in the nitrogen depth profiles. Additionally, the lattice parameters of

the pure γ phases and the nitrogen rich phases of iron and cobalt - FeN and CoN - in

the two possible cubic crystalline structures (NaCl-type and ZnS-type corresponding to

nitrogen on the octahedral sites and on the tetrahedral sites of the fcc lattice, respec-

tively) [138–142] are shown for stainless steel and CoCr. For comparison, further data

taken from literature (Christiansen et al. [122]) is indicated for stainless steel, where

thin austenitic stainless steel foils were nitrided in ammonia/hydrogen gas mixtures.

As can be seen for both types of alloys, a wide variation of the data points is present.

In addition, when fitting the data of stainless steel (figure 6.6 a), a linear dependency of

the lattice parameter on the nitrogen content is roughly visible. However, at nitrogen

contents lower than 0.17, deviations from this linear behavior are required to obtain

the lattice parameter of 3.6 × 10−10 m of the untreated base material. This indicates

that at low nitrogen contents a transition region exists where the expanded phase is

not fully developed. Only with higher nitrogen concentrations, the expanded austenite

will develop and the lattice parameter will increase. For CoCr (figure 6.6 b), a very

similar behavior is expected, however, the large variation of the data points does not

allow further conclusions, except that an additional influence of the process conditions

is dominating. Besides, it is rather difficult to produce samples with low nitrogen con-

tents.

The presented results confirm that equation (6.1) is only valid for a certain nitrogen

concentration range. Thus, a direct determination of the nitrogen content from X-ray

diffraction data is not unequivocal and should be avoided. Additionally, the calculated

concentrations represent average values as the broadened nature of the γN peaks indi-

cates a distribution of the nitrogen content.

64


0.0 0.2 0.4 0.6 0.8 1.0

3.6

3.8

4.0

4.2

4.4

4.6

FeN(ZnS-type)

FeN (NaCl-type)

-Fe

a) Stainless Steel

(111) (200) Christiansen et. al

Latti

ce P

aram

eter

[10-1

0 m]

Nitrogen Content [x]: FeNx

0.0 0.2 0.4 0.6 0.8 1.0

3.6

3.8

4.0

4.2

4.4CoN (ZnS-type)

CoN(NaCl-type)

-CoCr

b) CoCr

(111) (200) (111) (200)

Latti

ce P

aram

eter

[10-1

0 m]

Nitrogen Content [x]: CoNx

Figure 6.6: Lattice parameter of (111) and (200) reflections as function of nitrogen contentfor a) stainless steel (alloy 304) and b) CoCr (alloy HS188) for different PIII process condi-tions. The square symbols correspond to PIII with varying pulse voltages between 0 to -10 kVand a constant temperature of 420 ◦C for 2 h. Beforehand, the samples were preheated for1 h. In contrast, the spherical symbols correspond to PIII with -10 kV, a temperature of 390and 445 ◦C, respectively, and a process times of 2 h. Preheating was not performed. For com-parison, data taken from literature of gas nitriding of stainless steel is shown (Christiansenet al. [122]).The dotted lines indicate a linear relationship between nitrogen content and lattice constant as-suming the nitrogen atoms occupying the octahedral interstitial sites of the lattice (NaCl-type)as reported in literature.


The phase formation of stainless steel and CoCr alloys after nitrogen plasma immer-

sion ion implantation was investigated. X-ray diffraction after different PIII process

temperatures and annealing times leads to the following results:

• For both alloy classes, a characteristic expanded phase γN is formed at nitriding

temperatures up to 400 ◦C. Depending on the temperature and alloy, a lattice

expansion between 6 - 12 % is observed. At higher temperatures, the expanded

phase converses into CrxN and in a chromium-depleted phase. This chromium-

depleted phase consits of bcc Fe (ferrite) for stainless steel and an fcc CoNi phase

for CoCr. Excess nitrogen is supposed to agglomerate at the grain boundaries as

the ratio N:Cr > 1.

• The number of possible nitrogen sites in the fcc lattice can explain the limited

nitrogen surface concentrations between 30 - 40 at.% in the depth profiles (see

chapter 4). Occupying the octahedral sites in the fcc lattice with a maximum

65

6. PHASE FORMATION

of 2 nitrogen atoms per unit cell a concentration of 33 at% nitrogen is obtained.

Further nitrogen is much weaker bounded and can be easily released into the gas

phase.

• The decomposition rate of expanded austenite does not only depend on temper-

ature but also on time. Depending on the annealing temperature (400 - 500 ◦C),

decomposition rates between 4 h and less than 0.5 h are present. Higher temper-

atures correspond to a faster decomposition.

In addition, the formation of Cr2N is favored at lower temperatures and a slower

decomposition can be found for CoCr in contrast to stainless steel.

• Deviations from a time independent diffusion constant, thus from a parabolic

layer growth, at high temperatures were already mentioned in chapter 4.2: A

significant increase in the diffusion rate is observed for stainless steel, while a

sharp drop in nitrogen diffusivity is found in CoCr alloys. This correlates with

the decomposition of the expanded phase and the formation of bcc Fe (in stainless

steel) and fcc CoNi (in CoCr), respectively. The following model is proposed:

The decomposition of the expanded lattice starts from the surface towards the

bulk. Thus, the CrN/Cr2N containing matrix is near the surface region while the

expanded austenite is situated near the bulk material. Diffusion is increased in

the ferritic phase, whereas the fcc CoNi phase hinders the transport of nitrogen.

However, the layer thicknesses before and after annealing remain constant, even

after annealing temperatures more than 400 ◦C for 8 h. This is in accordance

with literature where an increase of the layer thickness was found after annealing

not until 500 ◦C and 20 h [143].

• It was shown that a linear correlation between nitrogen content and lattice pa-

rameter does not exist for low nitrogen concentrations. This suggests that at first,

a transition region is present before the “real” expanded austenite is formed, sepa-

rating expanded austenite from low concentration nitrogen contents. Thus, when

applying Vegard’s law these restrictions have to be considered.

Although a detailed investigation of stress and an anisotropic lattice expansion

was not performed, the annealing experiments showed that significant stress relax-

ations are present when cooling down the samples to room temperature. There-

fore, it has to be kept in mind that in general, X-ray diffraction data taken

after nitriding (and after cooling down the samples) gives information about the

conditions during nitriding to a limited extent only.

66

7 Surface Properties

The impact of low-energy, high-current density, nitrogen ion beams and plasmas on

the mechanical and electrochemical properties of stainless steel is well established and

several publications exist [65, 143–148]. In general, the insertion of nitrogen ions at

moderate temperatures (< 400 ◦C) leads to an improvement of hardness and wear resis-

tance whereas the corrosion properties remain unaffected. Thus, the following chapter

will focus on CoCr alloys.

As already mentioned, CoCr alloys have a wide application as biomaterial as they

exhibit excellent mechanical properties and a satisfying biocompatibility [2]. Never-

theless, the generation of nanoparticles by mechanical wear processes is still observed

in some total hip replacements after explantation [3]. Besides, fretting corrosion can

lead to the release of toxic ions such as cobalt, chromium and nickel [4, 5]. Therefore,

the aim of this chapter is not only to investigate the mechanical and tribological prop-

erties of nitrided CoCr alloys, but also the electrochemical behavior. Additionally, the

wear mechanisms in physiological solution will be investigated as the majority of the

experiments found in literature is either to study the wear mechanisms or the corro-

sion behavior independent from alternative influences. Although one publication on

improved tribological properties in sodium chloride solution exists [18], the interplay

between wear and corrosion together has not been addressed in detail yet.

The PIII treatments in this chapter were performed at a pulse voltage of -10 kV, pro-

cess temperatures between 340 and 555 ◦C and process times of 2 h. Afterwards, the

electrochemical surface properties were investigated with potentiodynamic polariza-

tion and electrochemical impedance spectroscopy (EIS) in Ringer solution containing

147.0 mM NaCl, 4.32 mM CaCl2 and 4.04 mM KCl - similar concentrations as their

occurrence in body fluids [91]. The potentiodynamic polarization scans started from

-200 mV relative to the free corrosion potential with a scan rate of 0.2 mV/s. The

cathodic branch of the polarization curve was used to determine the corrosion rate

using the Tafel slope [149]. EIS measurements were carried out over a frequency rage

from 10 kHz to 0.01 Hz. The amplitude of the sinusoidal signal was 10 mV.

Hardness was measured with a dynamic nanoindentation system equipped with a

67

7. SURFACE PROPERTIES

Berkovich tip at a load of 10 mN. The wear tests were performed using a conventional

oscillating ball-on-disc configuration with an alumina ball: ball diameter 4.76 mm, load

1 N, sliding distance 10 mm at a maximum speed of 10 cm/s. For medical applications,

this would correspond to the loading of a contact between a CoCr femoral head of a

total hip replacement and a corresponding ceramic acetabulum. The commonly used

medical combination with a CoCr ball as counter body was not possible to test as

CoCr balls are nearly impossible to procure. Furthermore, three-dimensional distribu-

tions of the von Mises stress below the contact area [150] were calculated using the

commercial software Elastica R©3.06. The calculations were performed for nitrided as

well as for untreated samples, both with and without lateral forces, in the latter case

assuming a generic friction coefficient of 0.20. Additional wear tests were performed to

compare the tribological behavior in air and in physiological solution in the same ex-

perimental set-up. The parameters were the following: load 1 N, speed 5 cm/s, radius

of circular wear tracks 5 mm, number of cycles 20000. The experiments were made

in air and in simulated body fluid (SBF). SBF was prepared according to Kokubo

et al. [151] containing 142.0 mM Na+, 5.0 mM K+, 1.5 mM Mg2+, 2.5 mM Ca2+,

147.8 mM Cl−, 4.2 mM HCO3−, 1.0 mM HPO42− and 0.5 mM SO42−. Analysis of the

samples’ morphology was done with scanning electron microscopy (SEM), their chemi-

cal composition was analyzed with energy-dispersive X-ray analysis (EDX) and GDOS.

In order to investigate the wear volume a laser profilometer was used. Additionally,

to determine the released ion species during the wear tests in SBF, the electrolyte was

analyzed with inductively coupled plasma optical emission spectroscopy (ICP).

7.1 Corrosion Behavior

Potentiodynamic polarization curves are shown in figure 7.1 a) for different PIII process

temperatures in comparison with the untreated specimen of CoCr alloy L605. One can

see that the samples treated at low temperatures (340 ◦C and 400 ◦C) remain passive

in the solution and their potentials are largely shifted relative to the untreated alloy.

Although their passive region is smaller, the passive current density is lower and similar

break through potentials are obtained (600 mV). However, PIII at elevated temper-

atures (555 ◦C) prevent the formation of a stable passive film and active dissolution

occurs straight after the corrosion potential. The respective corrosion rates calculated

by Tafel slope evaluation confirm these observations. Whereas relatively low corro-

sion rates (< 3 µm/year) are obtained at low PIII process temperatures, the corrosion

68

7.1. CORROSION BEHAVIOR

10-3

10-2

10-1

100

101

-250 0 250 500 750

Potential [mV]

Curr

ent[µ

A/c

m²]

untreated340 °C

400 °C

555 °C

10-3

10-2

10-1

100

101

-250 0 250 500 750

Potential [mV]

Curr

ent[µ

A/c

m²]

untreated340 °C

400 °C

555 °C

0

4

8

12

16

Corr

osio

nR

ate

[µm

/ye

ar]

PIII Process Temperature [°C]

untreated 340 400 5550

4

8

12

16

Corr

osio

nR

ate

[µm

/ye

ar]


untreated 340 400 555

a) b)

Figure 7.1: Potentiodynamic polarization curves (a) obtained for CoCr alloy L605 treated atthree different PIII process temperatures as well as for the untreated alloy in Ringer solution.The corresponding corrosion rates were determined by Tafel slope evaluation (b).

increases up to 17 µm/year at high temperatures. Nevertheless, the lowest corrosion

is obtained for the untreated alloy. In addition, similar tendencies are visible with

electrochemical impedance spectroscopy where the obtained charge transfer resistance

agrees well with the corrosion rates.

To get an insight into segregation effects near the surface, depth profiling was carried

out with GDOS. Figure 7.2 presents the respective depth profiles of cobalt, chromium

and oxygen of the specimens after the potentiodynamic measurements. The profiles of

the other elements present in the alloy do not differ from each other and are therefore

not shown. The untreated sample (a) exhibits the highest oxygen concentration of

48 at% near the surface in contrast to the untreated ones. In addition, a maximum in

the chromium concentration up to 43 at.% is visible in a depth of 10 nm in the un-

treated alloy but also in those samples nitrided at low temperatures (340 and 400 ◦C).

The sample treated at 555 ◦C, however, shows a much lower chromium content of only

about 5 at.% (in a depth of 10 nm). Besides, the cobalt content is slightly reduced.

The presented results confirm that the absence of chromium relates with the enhanced

corrosion rate. As a result, the formation of a passivating Cr2O3 oxide layer on the

surface is suppressed, compromising the corrosion resistance [152]. Further details on

the released ion species and the corrosion mechanisms will be discussed in section 7.3.

69


0 10 20 30 40 500

10

20

30

40

50

60

N

O

Cr

Co

a) untreated

Conce

ntratio

n[a

t.%

]

Depth [nm]

0 10 20 30 40 500

10

20

30

40

50

60

N

O

Cr

Co

b) 340 °C

Conce

ntratio

n[a

t.%

]

Depth [nm]

0 10 20 30 40 500

10

20

30

40

50

60

N

O

Cr

Co

c) 400 °C

Conce

ntratio

n[a

t.%

]Depth [nm]

0 10 20 30 40 500

10

20

30

40

50

60

N

O

Cr

Co

d) 555 °C

Conce

ntratio

n[a

t.%

]

Depth [nm]

Figure 7.2: Depth profiles of cobalt, chromium, oxygen and nitrogen in CoCr alloy L605measured with GDOS for the untreated alloy (a) and after nitrogen PIII at different temper-atures (nitrogen layer thicknesses: b) 0.8 µm, c) 1.3 µm and d) 2.8 µm).

7.2 Nano Indentation and Wear

The nanohardness of the surface layer is shown in figure 7.3. The hardness increases

from about 5 GPa (base material), towards 12 - 16 GPa at a load of 10 mN for CoCr

alloys L605 and SY21. The lower values at lower process temperatures are an artifact

of the hardness measurement as the 10 %-rule is violated, which means that the inden-

tation depth is larger than 10 % of the modified layer thickness [153]. Furthermore,

the relatively large error bars are related with the surface roughness which increases

70

7.2. NANO INDENTATION AND WEAR

Figure 7.3: Hardness ofCoCr alloys L605 and SY21as function of process tem-perature (PIII parameters:-10 kV, 2 h) The respectiveload was 10 mN.

0 200 300 400 500 6002

4

6

8

10

12

14

16

18

20

untreated

CoCr alloy L605 SY21

Har

dnes

s [G

Pa]


with increasing ion fluence and thus with increasing temperature (see chapter 4.1). In

the treatment temperature range from 400 to 550 ◦C, a constant hardness is obtained,

however different for the two alloys, whereas the base materials exhibit no differences

for these alloys. The strengthening mechanism can be ascribed to solid solution hard-

ening due to the high interstitial atomic content of nitrogen [55]. In addition, stacking

faults on the fcc sublattice may contribute to the strengthening. Furthermore, dif-

ferences in the alloy compositions could lead to variations in the elastic modulus and

yield strength (with the latter one closely related to the measured hardness [154]) and

therefore to different hardness values when comparing the two alloys. Beyond 550 ◦C,

the formation of CrN precipitates apparently leads to an equalization of the hardness

values.

The wear tests using an alumina ball as counter body lead to the following results

(figure 7.4): no dependence of the wear volume on the wear path - or number of cy-

cles - is observed for the untreated base material. However, the absolute values are

differing between 2.5 and 7.0 × 104 µm3/m where SY21 (a) shows the lowest value

and HS188 (b) the highest. After PIII treatment the differences between the alloys

nearly disappear and the absolute wear rates depend only on the PIII process temper-

ature. The layer thickness itself is high enough to support the mechanical load without

breakthrough during the experiment, except for HS188 implanted at 335 ◦C, where

the wear behavior of the non-implanted base material is observed at 4000 cycles and

beyond. For MP35N and HS188, a decrease of the wear rate with increasing process

temperature is observed, whereas a nearly inverse relation was found for SY21 with

71


0.0

2.0x104

4.0x104

6.0x104

0 2000 4000 6000 80000.0

2.0x103

4.0x103

4.0x104

6.0x104

0

1x103

2x103

2.0x104

4.0x104

a) CoCr SY21

untreated 335 °C 410 °C 440 °C 480 °C

c) CoCr HS188

Number of Wear Cycles

b) CoCr MP35N

Wea

r Vol

ume

[µm

3 /m]

Figure 7.4: Wear volume asa function of wear cyclesand PIII process temperaturefor CoCr alloy a) SY21, b)MP35N and c) HS188. Theload during each test was 1 N.Note the different scales andthe axis breaks.

the lowest process temperature leading to the lowest wear rate among the implanted

samples. The differences of wear volume due to temperature treatment are not very

significant above 400 ◦C. It has to be noted that no breakthrough of the layer is oc-

curring under these conditions, thus the difference and similarities in the wear rate

must be indicative of intrinsic properties of the nitrided surface layer. No differences

in the measured friction coefficient µ were observed with values of 0.50 - 0.55 present

across all measurements. Thus, the increase in the surface roughness during the PIII

treatment is still minor enough to avoid a strongly corrugated surface structure with

a subsequently higher friction coefficient. The most likely cause for the improved wear

rate is an increased cohesion or yield strength after the nitrogen insertion. Despite the

formation of an apparently identical expanded austenitic lattice with no influence of

the alloy on the nitrogen diffusivity, the intrinsic hardness and wear rate still depends

on the alloy composition.

An analysis of the wear track geometry and wear particles will be discussed in the next

72

7.2. NANO INDENTATION AND WEAR

section (chapter 7.3).

To have an insight into the stress distribution, calculations of the von-Mises stress -

a generalization of the stress observed in one-dimensional yield strength experiments

[155] - were performed. The obtained results are shown in figure 7.5. The two pan-

0a) Static (µ = 0) b) Dynamic (µ = 0.2)

0a) Static (µ = 0) b) Dynamic (µ = 0.2)

10 µm

10

µm

10 µm1

0µ

m0.70= 1.6 *

00.090.18

0.350.430.530.61

Stress[GPa]

0.26

1.10= 2.4 *

00.140.28

0.550.690.830.96

Stress[GPa]

0.41

Figure 7.5: Simulation of the von-Mises stress in 2D-cross-sections a) without (µ = 0)and b) with (µ = 0.2) lateral loading. The simulation shows the results of an alumina ball(diameter 4.76 mm) on CoCr assuming a yield strength δ = 0.45 GPa. A layer thickness of2 µm on top of the bulk material is assumed in both cases. Without lateral loading, a typicalindentation experiment with a ball indenter is simulated, whereas with lateral loading, thestress distribution during a wear test with a corresponding friction coefficient is shown. Thechange of the label of the stress scale from white to black indicates that the yield strength of0.45 GPa is exceeded.

els compare the loading configuration of 1 N under a) static (µ = 0) and b) dynamic

conditions with lateral load (µ = 0.2) (simulation of real wear experiment). The yield

strength δ of medical CoCr alloys is around 400 - 500 MPa and therefore the calculated

stress maximum is beyond the yield strength. As plastic deformation is not included in

the analytical calculations, which assume purely elastic behavior, a direct comparison

with the reality should not be made without keeping this limitation in mind. Neverthe-

less, a quite satisfactory agreement is found in literature in the case of stainless steel,

where the area of plastic deformation coincides with the area of stress exceeding the

yield strength [156].

As can be seen, static loading indicates a stress maximum located below the surface

layer within the base material. However, even a friction coefficient of 0.2 leads to a dis-

placement of the stress maximum towards the trailing edge of the contact area, while

the stress in the base material is reduced to values below the yield strength. For a

friction value of 0.50 - 0.55 encountered in the experiment, even larger deviations from

the static situation are present (not shown).

Comparing these results with the real wear experiments, an excellent adhesion of the

expanded austenite layer can be interfered from the data as no flaking or delamination

73


1 µm 5 µm

10 µm

2 µm 5 µm

5 µm

untreated

PIII 390 °C

PIII 570 °C

Wear in air Wear in SBF

a)

c)

e)

b)

d)

f)

1 µm 5 µm

10 µm

2 µm 5 µm

5 µm

untreated

PIII 390 °C

PIII 570 °C

Wear in air Wear in SBF

a)

c)

e)

b)

d)

f)

Figure 7.6: SEM micrographs of the wear tracks in air and in simulated body fluid (SBF) ofthe untreated specimens (a and b), the specimens nitrided at 390 ◦C (c and d) and the weartracks of the specimens nitrided at 570 ◦C (e and f) of CoCr alloy SY21.

was observed in the SEM micrographs.

7.3 Tribocorrosion

In the following, the wear mechanisms in physiological solution will be compared to

the wear behavior under dry conditions after PIII. The aim is to reflect the situation

of implants in medical environment.

The investigation of the wear tracks of CoCr alloy SY21 in air after the end of the

74

7.3. TRIBOCORROSION

experiment using SEM reveals a mixture of adhesion, abrasion and plastic deformation

for all specimens with much milder wear behavior for the PIII treated ones. Selected

wear tracks are shown in figure 7.6. Two groups of particles depending on the surface

treatment can be distinguished: spherical particles with a diameter of about 50 nm

(sometimes agglomerated to particles with a diameter of up to 1 µm) and needle-

shaped particles with a length of less than 1.5 µm. Whereas the spherical particles are

predominant in the wear tracks of untreated CoCr (figure 7.6 a and b) and nitrided

CoCr at high temperature (e and f, however, not visible in the magnification of panel f)

the needle-shaped particles characterize the specimens nitrided at low temperatures (c

and d). As shown by EDX measurements they both have an increased oxygen content

in comparison with the surrounding surface. The occurrence of nanosized spherical and

needle-shaped wear particles of untreated cobalt-based alloys has been already reported

in literature. According to the work of Buscher et al. [157] the spherical particles re-

sult from torn-off nanocrystals while the needle-shaped ones are generated by fractured

ε-martensite (hcp martensite). As no diffraction lines according to ε-martensite are

visible in the X-ray diffraction patterns, the ε-martensite grains are apparently too

small to be detected with X-ray diffraction. The occurrence of the expanded austenitic

lattice after nitrogen insertion at moderate temperatures is obviously highly correlated

with the occurrence of ε-martensite in the samples of this work. In contrast, the CrN

precipitates observed at higher temperatures do not modify the underlying wear pro-

cess. Thus, an intentional aging with a transformation of the initial fcc lattice can be

excluded, while differences in the stacking fault energy for the base material and the

expanded austenite structure may explain the modified particle formation [158].

However, after PIII treatment at high temperature the wear regime in SBF (figure 7.6 f)

changes completely to a much more aggressive wear compared to the sample in air. The

surface suffers strongly from fretting which is revealed by cracks and delamination. A

synergistic interplay of wear and corrosion seems to play a decisive role after nitriding

at elevated temperatures.

The respective normalized wear volume is given in figure 7.7 as a function of PIII pro-

cess temperature. It decreases in air from 39 × 104 µm3/m for the untreated material

to values lower than 0.5 × 103 µm3/m for the nitrided material with only a slight

dependency on the process temperature. The change between spherical and needle-like

particles is not reflected in the wear rate. In SBF, the wear resistance improves signifi-

cantly at low nitriding temperatures by a factor of ten to 0.4 × 103 µm3/m. However,

at process temperatures higher than 400 ◦C the wear rate increases again, even exceed-

ing the value of the untreated sample in SBF. These results are in agreement with the

75


0 350 400 450 500 550 6000.0

2.0x103

4.0x103

6.0x103

2.0x104

4.0x104

6.0x104

untreated

Wear test in air Wear test in simulated body fluid

Wea

r Vol

ume

[µm

3 /m]


Figure 7.7: Wear volume asa function of PIII processtemperature of wear tests inair and in SBF of CoCr al-loy SY21. The respective loadduring the wear test was 1 N.

SEM micrographs and suggest that a combination of corrosion and wear, i.e. fretting

corrosion, leads to an elevated wear rate in SBF compared to the wear experiments in

air.

To ascertain the corrosion rate, the elemental concentrations of the metals present in

the alloy in the SBF were obtained by ICP measurements after the wear tests, as pre-

sented in figure 7.8 a). From the main alloying elements, only chromium and cobalt

0.0

0.2

0.4

0.6

0.8

1.0

1.2 a)

570530390 405untreated PIII Process Temperature [°C]

Ion

Con

cent

ratio

n [m

g/l]

Co (x 0.1) Cr Ni

0

1x1010

2x1010

3x1010

4x1010

5x1010

6x1010b)

PIII Process Temperature [°C]570530405390untreated

Num

ber o

f Wea

r Par

ticle

s

Wear in SBF

Figure 7.8: Comparison of a) metal ion release of alloy SY21 during wear tests in SBF withb) number of wear particles after nitrogen PIII at different temperatures. Note that the cobaltconcentrations were multiplied by a factor of 0.1 for clarity.

were investigated, as the tolerance levels of molybdenum are quite high for humans

[159]. Additionally, nickel was investigated despite a content below 1 at.% in the inves-

tigated alloy, as a very high toxicity is known for this element [160]. Different behavior

76

7.3. TRIBOCORROSION

of the untreated alloy and the nitrided alloys at different temperatures are found. While

the amount of cobalt shows a slight increase of a factor of 2.5 mg/l from 0.3 for the

untreated material to 0.75 mg/l for the nitrided material at 390 ◦C, despite a lower

wear rate for this sample, for even higher temperatures a strong rise of cobalt up to

11.3 mg/l is observed. When comparing the concentrations of chromium and nickel

only small variations are seen between the untreated and treated alloys. All concentra-

tions are below 0.1 mg/l with the lowest values for the sample nitrided at 405 ◦C. When

the increase of the surface roughness with treatment temperature is taken into account,

the ion release rates normalized to the actual surface will be even lower than indicated

as a rougher surface provides a greater surface area for metal ion release. However,

comparing the ratio of surface roughness increase with ion release, it is obvious that

the high ion concentrations of cobalt in solution cannot only be due to the roughened

surface.

The high cobalt release at different implantation temperatures may be attributed to

the stability of the nitrogen implanted layers formed under different implantation con-

ditions as already indicated in the potentiodynamic polarization measurements. At low

temperatures the expanded austenitic structure with nitrogen in solid solution in gen-

eral maintains the passivating nature of the original CoCr surface. However, a stronger

affinity of nitrogen to chromium than to cobalt could lead to a partial weakening of

the original chemical bonding within the metallic alloy, similar to the situation found

for nitriding of austenitic stainless steel [122]. Furthermore, elevated temperatures

which lead to the formation of chromium nitrides, result in a depletion of chromium

in solid solution prohibiting the formation of a passivating Cr2O3 surface oxide and

by allowing contact corrosion between the matrix and the precipitates. For compar-

ison, in figure 7.8 b) the number of wear particles in SBF is calculated assuming an

average particle size of 50 nm and the wear volume in SBF already presented in fig-

ure 7.7. Although the untreated sample exhibits the best corrosion resistance (i.e.

the lowest Co ion release rate) the wear rate is much higher than those of the PIII

nitrided samples. Therefore, on the one hand the nitrogen insertion leads to either

an expanded austenitic structure or to the formation of chromium nitride precipitates

at elevated temperatures which both result in improved tribological properties in air.

On the other hand, the passivating behavior of the untreated material disappears with

increasing process temperatures and leads to a continuous, selective increase in the

cobalt ion release. Consequently, fretting corrosion occurs with increasing PIII tem-

perature, while the wear rate actually shows a minimum at intermediate PIII process

temperatures.

77



The results of this chapter can be summarized as follows:

• The insertion of nitrogen affects the electrochemical properties of CoCr. However,

whereas PIII at low temperatures (< 400 ◦C) only leads to a slight reduction of

corrosion resistance, nitriding at high temperatures leads to an active dissolution

of the surface and an increase of the corrosion rate up to a factor of six. The

loss of passivity is attributed to changes in the lattice structure due to nitrogen

treatment. Similar to austenitic stainless steel, the precipitation of CrN at high

PIII process temperatures, prevents the formation of a passivating Cr2O3 oxide

layer and corrosion can take place. This is confirmed by GDOS analysis where a

depletion of chromium near the surface is found for the PIII treated samples at

high temperatures.

• The dominating released ion species during exposure to physiological solution is

cobalt. Any nitrogen implantation results in surfaces with higher levels of cobalt

release than the untreated surfaces, whereas no influence of the PIII treatment

on the nickel or chromium dissolution is observed. This may be explained by the

nature of the implanted layer. As the nitrogen atoms on the interstitial sites of

the fcc lattice have a stronger bond with chromium than with cobalt, cobalt can

be easily released.

• Hard and wear resistant surface layers are obtained: The hardness increases up

to a factor of three (from 5 to 16 GPa at a load of 10 mN) while the wear rate is

roughly decreased by a factor of up ten, depending on the properties of the base

material (from about 6 × 104 to 0.6 × 104 µm3/m at a load of 1 N). Initial dif-

ferences in the mechanical properties of the different CoCr alloys are eliminated

to a large extent by the surface modification. This indicates that the dominating

factor is the nitrogen insertion, either as an expanded austenitic structure de-

rived from the fcc phase of the CoCr alloys or from the CrN precipitates arising

at higher temperatures. The PIII process temperature itself has only a minor

influence in determining the layer thickness by thermally activated diffusion.

• Calculations of the stress distribution yield in maximum stress values up to 1 GPa

within the bulk material and the surface layer. These values are beyond the yield

strength of about 400 - 500 MPa and therefore, cold working of the material may

take place. No layer delamination or interface failure between the modified surface

78


and the bulk was observed indicating an excellent cohesion of the implanted

surface layer.

• The friction coefficient is nearly unchanged (0.50 - 0.55) after nitrogen insertion.

This indicates that the surface roughening due to the ion bombardment during

PIII does not increase the friction coefficient. At the same time, an increased

cohesion is observed after nitriding as the energy dissipation is identical. The

wear mechanism itself is abrasive wear of the softer CoCr by the alumina counter

body. Adhesive wear or cold welding was not observed in the present experiments.

• Comparing the results of tribocorrosion after nitrogen PIII in simulated body

fluid to the wear behavior in air, a dependency of the wear mechanism on the

PIII process temperature is found. Whereas untreated CoCr suffers from strong

abrasive wear, the nitrogen implantation at low temperatures leads to improved

tribological properties with very small abrasive wear and only a minor contri-

bution of fretting corrosion. After PIII at high temperatures beyond 400 ◦C,

fretting corrosion dominates and results in a strong increase in the wear rate and

the excessive release of cobalt ions. The presented results favor PIII treatments

at temperatures below 400 ◦C to achieve an optimal interplay between low wear

rates and high corrosion resistance. However, an additional balancing of low par-

ticle release rates or low cobalt ion release rates has to be performed, depending

on the specific application.

79

8 Summary and Conclusions

In this work, nitrogen plasma immersion ion implantation of austenitic base alloys was

performed and the influence of their alloy composition on the resulting physical proper-

ties was investigated. The main focus was on the characterization of nitrogen diffusion

and subsequent layer formation, phase formation and the resulting surface properties,

i.e the mechanical and electrochemical properties.

Two groups of alloys - Co-Cr-Ni and Fe-Cr-Ni - with similar grain sizes between 10 and

50 µm were chosen whose austenitic γ phase was stable over the whole temperature

range. Thus, spurious influences of either varying grain sizes or phase transitions of

the base material were avoided.

It was shown that CoCr alloys, similar to stainless steel, can be efficiently nitrided

within a few hours using PIII at temperatures between 300 and 600 ◦C. The diffusion

process is characterized by thermally activated diffusion resulting in nitrogen enriched

layers with thicknesses up to several micrometers. The interstitial diffusion of nitro-

gen is nearly independent of the alloy composition, with similar activation energies

(0.9 - 1.1 eV) and only small differences of the exponential prefactors between steel

and CoCr alloys. An unusual plateau-like diffusion profile at 25 - 40 at.% is found for

both groups with nitrogen supply essential for diffusion. Trapping models in literature

underestimate the energy necessary to remove nitrogen again from traps, as proven in

annealing experiments without further nitrogen flux performed in this work. The traps

themselves are related to interstitial sites. Chromium atoms seem to play a minor role

for diffusion as the retained nitrogen concentration is much higher than the chromium

content of the base material and independent thereof (chromium content before nitrid-

ing is 20 - 30 at.%).

In contrast, a strong dependency of nitrogen incorporation during PIII on the alloy

composition, properties of surface oxides and on further effects originating from inter-

actions of the surface with gas, plasma and energetic ions could be found. Thus, a

detailed qualitative model was presented, delineating different processes at the surface:

dissociative radiation-enhanced adsorption, transport from the surface through a bar-

rier, consisting primarily of a surface oxide, and the diffusion towards the bulk. It was

81

8. SUMMARY AND CONCLUSIONS

shown that the rate-limiting factor during PIII is the diffusion process in combination

with the maximum surface concentration of nitrogen in the layer and not the nitrogen

supply. Therefore, the faster diffusion of nitrogen in stainless steel in comparison to

CoCr allows a larger uptake of nitrogen. In addition, not only implantation of energetic

ions but also the nitrogen background pressure and adsorption of nitrogen molecules

during the pulse pauses play a decisive role for the nitrogen uptake. The impinging

ions during the high voltage pulses dissociate these molecules to atoms which diffuse

into the material. An oxide layer does not affect the uptake of nitrogen during PIII, as

it will be sputtered away - at most it will delay the uptake.

In contrast, for interactions of CoCr and stainless steel with a plasma only, a strong cor-

relation of the diffusivity with the chemical composition of the surface and the nature

of the surface oxide was identified. For CoCr alloys, the rate limiting step is mainly

the surface adsorption, while the oxide layer is a perfect barrier only in the case of

austenitic stainless steel. These results imply that in the case of CoCr, nitrogen PIII

could be replaced by a less intricate plasma nitriding process (e.g. as developed in this

work) when the surface adsorption is optimized. In addition, plasma nitriding would

avoid the strong roughening, which is desired for some applications where smooth sur-

faces are necessary.

For both classes of alloys, nitrogen PIII at temperatures up to 400 ◦C leads to an

expansion of the fcc lattice between 6 - 12 %. At temperatures beyond 400 ◦C, the

vacancy assisted diffusion of Cr atoms leads to a decomposition of the metastable ex-

panded phase into chromium nitride precipitates and a chromium-depleted phase. For

steel, ferrite was observed with a fast nitrogen diffusion inside the ferrite, whereas an

austenitic CoNi phase is remaining for CoCr alloys, where no observable nitrogen diffu-

sion is occurring. However, additional nitrogen is either present at the grain boundaries

of the precipitates or within the matrix as the N:Cr ratio is always larger than one.

Additionally, the precipitation process of chromium nitrides does not only depend on

temperature but also on time, as a time delay is observed which is not present for the

nitriding process itself. This leads to the formation of a two-layer system consisting of

CrN/γN/γ-bulk with different growth rates for the two layers.

The formation of an expanded austenitic lattice is observed for both alloy groups. For

stainless steel, the expansion itself is smaller than predicted from the lattice constants

of FeN when taking the actual nitrogen content into account. However, for the initial

low nitrogen content region, a threshold seems to be identifiable, beyond which the ex-

panded lattice characterized by the unusual diffusion behavior is observed. For CoCr

and stainless steel, no clear correlation between the expansion and the nitrogen content

82

is observable, indicating that detailed process conditions, including the stress created

by the implanted atoms, are more influential than the nitrogen content alone. It could

be shown that Vegard’s law that up to now is often used in literature to calculate the

nitrogen content from the lattice expansion, should be avoided.

The limited maximum surface concentration in the nitrogen enriched layers (∼ 30 -

40 at.%) was associated with a limited number of possible nitrogen sites at the octa-

hedral sites of the fcc lattice. Further implanted nitrogen is much more weakly bound

and can therefore easily be released.

The possibilities and restrictions of PIII nitriding of CoCr alloys for practical appli-

cations - especially for medical application - were presented. It was shown that the

nitrogen insertion leads to solid solution hardening, resulting in a significant improve-

ment of hardness of a factor of three and an increase of wear resistance of up to factor

of ten. Hardness values and friction coefficients are again similar when comparing

the alloys after nitriding, while differences in the wear rate are observed. For these

hardening-, and the priorly discussed diffusion processes as well, an average potential

of the metallic solid-solution should be responsible, whereas abrasive wear depends on

the yield strength, i.e. cohesive energy and stacking fault energy, which are crucially

depending on specific interatomic potentials. Thus, the selective dependency of the

wear rate on the composition can be understood. Depending on the process temper-

ature, different kinds of wear particles with an increased oxygen concentration were

found: spherical particles (diameter ≈ 50 nm, sometimes agglomerated to particles

with a diameter of up to 1 µm) and needle shaped particles (length ≈ 1.5 µm).

However, the surface treated alloys suffer from degradation of the corrosion resistance.

While in steel, the corrosion resistance changes between excellent and poor when the

transition towards CrxN formation is reached, an impairment of corrosion resistance

is observed across the whole PIII temperature range in CoCr alloys that additionally

increases with increasing temperature. The selective release of cobalt as the only ion

species into the surrounding solution was attributed to the nature of the nitrogen en-

riched layer. The insertion of nitrogen degrades the stability of the Co-Cr-Ni compound

by weakening the cohesion of cobalt atoms, leading to a much easier release thereof.

Detailed investigations on the electronic band structure would be necessary to eluci-

date this process. At even higher process temperatures, similar effects as present in

stainless steel develop. The precipitation of CrxN immobilizes the chromium, no longer

allowing the formation of a passivating Cr2O3 oxide layer on the surface, which results

in a deterioration of the corrosion resistance. Experiments to test the biocompatibility

of the PIII treated CoCr alloys are still in progress. However, first results with cells

83

8. SUMMARY AND CONCLUSIONS

grown on the nitrided CoCr surfaces are on the same lines of reasoning: An elevated

cobalt release is found for the PIII treated alloys at high temperatures, which impairs

the vitality of the cells. In contrast, after surface treatment at low temperatures, only

a minor reduction of the vitality is found.

As most publications in literature deal with either studying the metal ion release or

the wear behavior, an additional focus of this work was to combine both processes and

to study the interplay of corrosion and wear together. Thus, the investigations of the

wear mechanisms in physiological solution showed that the wear changes from strong

abrasive wear of the untreated alloys to slight abrasive wear and minor fretting corro-

sion after PIII at temperatures below 400 ◦C. Finally, after PIII at high temperatures,

slight abrasive wear is found, however, it is accompanied by strong fretting corrosion.

Therefore, depending on the application, a compromise for the PIII treatment temper-

ature must be found as the minimum wear particle release rate does not correspond to

the minimum cobalt ion release rate. The still open question is whether cellular dam-

age by nanoparticle uptake and cobalt ion uptake can be reduced after PIII treatment

or not.

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94

List of Figures

2.1 Development of the concentration profiles during sputtering . . . . . . . . . . 16

3.1 Fcc lattice . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

3.2 Schematic experimental setup of the UHV chamber . . . . . . . . . . . . . . . 23

3.3 Sketch of the ECR plasma source . . . . . . . . . . . . . . . . . . . . . . . . . 24

3.4 Geometry of an X-ray diffractometer in Bragg-Brentano geometry . . . . . . 27

4.1 SIMS profiles of alloy HS188 . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32

4.2 Comparison of nitrogen profiles measured by GDOS and SIMS . . . . . . . . 33

4.3 Typical SIMS depth profiles of CoCr and stainless steel . . . . . . . . . . . . 34

4.4 Cross-section of implanted CoCr alloy . . . . . . . . . . . . . . . . . . . . . . 35

4.5 Comparison of roughness evolution as function of fluence . . . . . . . . . . . . 35

4.6 Arrhenius plot of resulting diffusion coefficients . . . . . . . . . . . . . . . . . 36

4.7 Temporal development of the layer thickness of CoCr . . . . . . . . . . . . . . 37

4.8 Temporal evolution of the layer thickness at low and high process temperatures 38

5.1 Ratio of retained-to-incident fluence . . . . . . . . . . . . . . . . . . . . . . . 43

5.2 Variation of nitrogen background pressure . . . . . . . . . . . . . . . . . . . . 44

5.3 Potential energy diagram as function of distance from the surface . . . . . . . 46

5.4 SIMS nitrogen depth profiles after PIII, GN and PN . . . . . . . . . . . . . . 47

5.5 Arrhenius plot of resulting diffusion coefficients after PIII and PN . . . . . . 47

5.6 SIMS nitrogen depth profiles after PIII and PN with native and thermal oxide

layers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49

5.7 SIMS depth profiles of the near surface-regions with thermal and native oxide

layers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50

5.8 Depth profiles after PN with preliminary sputtering with argon and nitrogen 51

5.9 Overview of processes during GN, PN and PIII . . . . . . . . . . . . . . . . . 53

6.1 XRD patterns after PIII at different temperatures . . . . . . . . . . . . . . . 57

6.2 Octahedral interstitial sites of the fcc lattice . . . . . . . . . . . . . . . . . . . 58

6.3 XRD patterns during annealing . . . . . . . . . . . . . . . . . . . . . . . . . . 59

6.4 XRD patterns after PIII at different process times . . . . . . . . . . . . . . . 61

95

List of Figures

6.5 Schematic development of phase formation and diffusion . . . . . . . . . . . . 63

6.6 Lattice parameter as function of nitrogen content . . . . . . . . . . . . . . . . 65

7.1 Potentiodynamic polarization curves and corrosion rates . . . . . . . . . . . . 69

7.2 Depth profiles of cobalt, chromium, oxygen and nitrogen after potentiodynamic

measurements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70

7.3 Hardness . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 71

7.4 Wear . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

7.5 Simulation of the von-Mises stress . . . . . . . . . . . . . . . . . . . . . . . . 73

7.6 SEM micrographs of the wear tracks . . . . . . . . . . . . . . . . . . . . . . . 74

7.7 Wear in air and in SBF . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 76

7.8 Comparison of metal ion release with number of wear particles . . . . . . . . 76

96

List of Tables

3.1 Chemical compositions of investigated alloys . . . . . . . . . . . . . . . . . . . 21

97

Acknowledgements

First of all, I want to acknowledge Prof. Dr. Bernd Rauschenbach for not only giving me

the opportunity to work on my thesis at the IOM, but also on his constructive advice and

support during the last years.

I want to thank PD Dr. Stephan Mandl for being a great supervisor, for his strong support

and all the extensive discussions. Without his advice and guidance this thesis would not have

been possible.

A lot of thanks go to Dr. Darina Manova for all the scientific and moral support. I am very

grateful for her encouragement and help during the last years.

Dr. Jurgen W. Gerlach is greatly acknowledged for performing several XPS and XRD inves-

tigations as well as for constructive comments and suggestions. At this point, I also want to

thank Moharram Abd El Khair and Artjom Bergmann for XRD measurements.

I wish to thank Dietmar Hirsch, Dr. Christian Patzig and Andrea Prager for SEM and ICP-

OES, respectively.

Some investigations would not have been possible without support from groups outside the

IOM Leipzig: PD Dr. Harm Wulff and Dr. Marion Quaas (Universitat Greifswald) are ac-

knowledged for X-ray diffraction measurements with annealing. I also want to thank Dr.

Carsten Blawert (GKSS Forschungszentrum Geesthacht) for the electrochemical characteri-

zation with potentiodynamic polarization scans and electrochemical impedance spectroscopy.

I am grateful to Ingrid Herold, Petra Hertel and Katharina Schulze for their kind support in

the lab and for assistance.

Working was not only of scientific profit but also a great pleasure due to the good atmosphere.

I would like to thank Christian, Yvonne, Antje, Inga, Marisa, Susi, Johannes, Manu, Andre,

Jens and Eric.

Finally, a very special thanks goes to my family, to my sisters Theresa and Katharina, to

Bjorn and to my parents for their unresented support and encouragement.

Selbststandigkeitserklarung

Hiermit versichere ich, dass die vorliegende Arbeit ohne unzulassige Hilfe und ohne Benutzung

anderer als der angegebenen Hilfsmittel angefertigt und dass die aus fremden Quellen di-

rekt oder indirekt ubernommenen Gedanken in der Arbeit als solche kenntlich gemacht wur-

den. Alle Personen, von denen bei der Auswahl und Auswertung des Materials sowie bei

der Herstellung des Manuskripts Unterstutzungsleistungen erhalten wurden, sind namentlich

genannt. Außer den in der Arbeit Genannten waren keine weiteren Personen bei der geistigen

Herstellung der vorliegenden Arbeit beteiligt. Insbesondere haben keine Personen von dem

Bewerber oder in seinem Auftrag unmittelbar oder mittelbar geldwerte Leistungen fur Ar-

beiten erhalten, die im Zusammenhang mit dem Inhalt der vorgelegten Dissertation stehen.

Es wird weiterhin versichert, dass die vorgelegte Arbeit weder im Inland noch im Ausland

in gleicher oder in ahnlicher Form einer anderen Prufungsbehorde zum Zwecke einer Pro-

motion oder eines anderen Prufungsverfahrens vorgelegt und in ihrer Gesamtheit noch nicht

veroffentlicht wurde. Es haben keine fruheren erfolglosen Promotionsversuche stattgefunden.

Documents

Diffusion Behavior and Phase Formation for Ion Implanted Austenitic