Surface Modification of Molds and Acessories For The Glass Industry
Surface Modification of Molds and Acessories For The Glass Industry
Surface Modification of Molds and Acessories For The Glass Industry
Doctor of Philosophy Thesis in Mechanical Engineering, Production Technology branch, supervised by Professor Altino de Jesus
Roque Loureiro and Professor Albano Augusto Cavaleiro Rodrigues de Carvalho and submitted to the Mechanical Engineering
Department of the Faculty of Sciences and Technology of the University of Coimbra.
2014
IMAGEM
Doctor of Philosophy Thesis in Mechanical Engineering, Production Technology branch, supervised by Professor Altino de Jesus
Roque Loureiro and Professor Albano Augusto Cavaleiro Rodrigues de Carvalho and submitted to the Mechanical Engineering
Department of the Faculty of Sciences and Technology of the University of Coimbra.
2014
Bolsa de Doutoramento (SFRH/BD/68740/2010)
Acknowledgments
Acknowledgments
Many people have contributed and encouraged me along all the years of this research, I thank
all of them, and I hope I have gathered a bit of their best.
Above all, I want to express my sincere gratitude to my supervisors: Albano Cavaleiro and
Altino Loureiro, for their scientific guidance, commitment and availability during the
execution of this thesis. Their bright mind and wisdom brought to me many insights and
knowledge that was the key success for this thesis. I’m also extremely indebted to them for all
the opportunities I had of involvement with various enterprises and scientific research centers
as well as of visiting several wonderful places in the world in the framework of conferences
and collaborations.
I have had the pleasure of working with and learning from a huge number of other professors
and researchers in different countries. I would like to specially thank professors Amilcar
Ramalho, Tomas Polcar, Josep Guilemany and Jerzy Morgiel for all the collaboration,
scientific discussion and open access to their laboratories.
Thanks are also due to INTERMOLDE company for the support and collaboration in the
initial stage of the thesis. Thanks are also extended to TEandM Enterprise.
It is essential to emphasize the crucial role of IPN (Instituto Pedro Nunes) on my work and
results, by making available laboratorial equipment for the development and characterization
of the samples I produced.
Special thanks to my friends and research colleagues from ECAT, Surface Engineering and
Nanomaterials and Micromanufacturing groups of CEMUC. I have spent marvelous moments
in their companionship. These thanks are extended to all the young and senior members from
i
Acknowledgments
the different international groups (Barcelona, Prague, Southampton and Guangxi (China)) I
visited, who welcomed me and helped me during my training periods there.
I would like to acknowledge the funding provided by the FCT (Fundação para a Ciência e
Tecnologia), through SFRH/BD/68740/2010 fellowship and the projects AUTOMATIN
“5380”, PLUNGETEC “13545” and PTDC/EME-TME/122116/2010.
Finally, I would like to thank my parents, grandparents and sister who permanently gave me
their support during this hard but reward time. Without their patience and good advices, I
would not have been able to take advantage of the opportunity I was given to improve my
education. This thesis is dedicated to them. Last but not the least, I thank you Cláudia for your
daily personal support and encouragement in all good and bad times.
ii
Abstract
Abstract
Coatings are frequently applied on molds and accessories for the glass industry in
order to restrict surface degradation such as oxidation, corrosion, abrasion and wear of the
structural material, thereby decreasing the maintenance costs and increasing the lifetime and
performance of the components. However, in order to obtain accurate lifetime expectancies
and performance of the coatings it is necessary to have a complete reliable understanding of
their properties.
This thesis is on the improvement of the surface properties and integrity of molds, in
order to increase their durability, through the application of different types of coatings. Two
methodologies were followed to reach such demands: (i) optimization of coatings currently
used in molds surface protection (Ni-based alloys deposited by Plasma Transferred Arc -
PTA); (ii) synthesis and characterization of new coatings with improved functionalities,
deposited by emergent deposition processes such as APS - Atmospheric Plasma Spraying
(effect of nanostructured ZrO2 additions on Ni-based alloy coatings) and DCRMS - Direct
Current Reactive Magnetron Sputtering (influence of V additions on the properties of
TiSi(V)N thin films).
The dilution of the substrate in PTA process was shown to strongly influence the
structure and, consequently, the hardness, the oxidation resistance and the tribological
behavior of the coatings; with increasing dilution, a detrimental effect on these properties was
observed due to the incorporation of base material. However, in relation to the tribological
behavior, a beneficial effect at high temperature was demonstrated due to the fast formation of
oxide layers which protect the coating surface against wear. The post-weld heat treatment
performed at coatings reduced the hardness of the partially melted and heat affected zones
without affecting the coatings hardness; whereas the coatings hardness and wear resistance
was improved with annealing treatment. Thus, the best performing coating could only be
achieved by a proper selection of the deposition conditions, in order to get the best
compromise between mechanical properties, high temperature oxidation behavior and wear
resistance of the coatings.
The impact promoted by nanostructured ZrO2 additions on the microstructure of a Ni-
based alloy depended on the way how the coatings were deposited by APS process, using: (i)
nanostructured ZrO2 and Ni-alloy powders previously mixed by mechanical alloying or (ii)
iii
Abstract
the same powders supplied separately. A homogeneous and compact microstructure with
small zirconia particles evenly distributed in the matrix was achieved in the first case, while a
porous microstructure, full of semi-melted Ni powders with large particles of ZrO2 entrapped
in their boundaries, suggesting a brittle behavior, was deposited in the second. In both cases
the hardness and wear behavior of ZrO2 rich coatings were improved in relation to the Ni-
based alloy. The coatings deposited from mechanically alloyed powders revealed to be much
more tribologically performing due to their compact structure and even distribution of
zirconia particles. All the APS coatings showed higher hardness values than the Ni-based
coatings deposited by PTA; however, their micro-abrasion resistance was worse, due to the
lack of cohesion between the powders.
The analysis by XRD of the structure of V rich TiSi(V)N coatings, deposited by
DCRMS, revealed that V incorporation in the TiSiN system shifted the peaks to higher
angles, indicating the formation of a substitutional solid solution based on TiN phase, where
Ti atoms are replaced by the smaller V ones. On the other hand, by similar reason, XRD of
TiSiN films revealed that a nanocomposite structure consisting of TiN grains enrobed by a Si-
N matrix was not formed. In fact, with Si addition a shift of the diffraction peaks of TiN phase
to higher angles was observed which, in combination with similar compressive residual
stresses, also supported the formation of a substitutional solid solution. V additions showed to
successfully improve the hardness and tribological behavior of TiSi(V)N films, as a result of
the substitutional solid solution formation. On the other hand, the formation of the V2O5 phase
during the sliding contact acts as a lubricious tribo-film, protecting the coating against wear.
The addition of Si or V strongly influenced in opposite directions the oxidation resistance of
the coatings. Si incorporations in TiN increased significantly the oxidation resistance of the
films, whereas the opposite occurred in V containing coatings. In this latter case, the rapid V
ions out-diffusion through the oxide scale inhibited the formation of a continuous protective
silicon oxide layer, which is responsible for the excellent oxidation behavior of TiSiN films.
V rich coatings showed lower oxidation resistance than PTA-deposited Ni-based coatings, but
superior hardness values and better tribological behavior was found in relation to both PTA
and APS deposited coatings.
Keywords: Glass Industry, Molds surface protection, Ni-based thick coatings, TiSi(V)N thin
films, self-lubricant coatings.
iv
Resumo
Resumo
v
* PTA – Do inglês – Plasma Transferred Arc
** APS – Do inglês – Atmospheric Plasma Spraying
*** PVD – Do inglês – Physical Vapor Deposition
Resumo
camada de óxido formada, que inibiu a formação de uma camada contínua e protectora de
óxido de silício, contrariamente ao que acontece nos revestimentos TiSiN. Os filmes ricos em
V apresentaram menor resistência à oxidação do que os revestimentos espessos depositados
por PTA; no entanto, a sua dureza e resistência ao desgaste foram bastante superiores à dos
revestimentos espessos depositados por PTA ou APS.
vii
* PTA – Do inglês – Plasma Transferred Arc
** APS – Do inglês – Atmospheric Plasma Spraying
*** PVD – Do inglês – Physical Vapor Deposition
Index
Index
Acknowledgments ....................................................................................................................... i
Abstract ......................................................................................................................................iii
Resumo ....................................................................................................................................... v
Index .......................................................................................................................................... ix
Nomenclature............................................................................................................................ xv
1. Introduction ......................................................................................................................... 1
3. Effect of the arc current variation on the properties of Ni-based coatings deposited by
PTA process .............................................................................................................................. 21
ix
Index
7. References......................................................................................................................... 59
Annex A ………………………….…………………………………...………………..…… 71
Annex B …………………………………………………………………………………….. 87
Annex C …………………………………………………………………………………….. 97
Annex D …………………………………………………………………………………… 113
Annex E ………………………………………………………………………………….… 123
Annex F ………………………………………………………………………………….… 145
Annex G …………………………………………………………………………………… 159
Annex H …………………………………………………………………………………… 171
Annex I ………………………………………………………………………………..…… 187
x
List of Figures
List of Figures
xi
List of Figures
xii
List of Tables
List of Tables
Table 2.1 - Comparison of the characteristics of several hardfacing and thermal spraying
processes [9, 55-60]. ................................................................................................................. 15
Table 2.2 - Typical sequential tasks used in molds production. ............................................... 17
Table 5.1 - Coatings deposited and their nomenclature. .......................................................... 40
xiii
Nomenclature
Nomenclature
xv
Chapter I – Introduction
Chapter I
1. Introduction
The necessity for higher performance and increased efficiency of components working
in extreme harsh conditions has been met with the use of more advanced materials. Coatings
have been developed over the last decades as a potential solution for these applications due to
their superior properties when compared to bulk materials, allowing extending the
performance and lifetime of components. This was the case of the glass industry, where the
components are actually coated with hard and wear resistant materials in order to resist to the
direct contact with the melted glass which promotes very severe abrasion, corrosion and
fatigue at high temperature. The current European glass container market is approximately 30-
40 billions of containers. Due to globalization, there has been a growing demand for glass
containers production at reduced costs, with increasing shape complexity, putting an
increasing pressure on glass producers who are seeking molds which allow increasing
production rates with higher quality. In these very harsh conditions, further degradation of the
mold surfaces occurs due to more intense thermal cycles leading to higher mechanical strains
and plastic deformation with the consequent fatigue, oxidation and wear problems.
Furthermore, other common problems such as worn-out geometries or surface and base
material cracks, often lead to the breakdown of the coating after few working cycles. The
combined effect of the harsh conditions with a faulty design or a defective material,
associated with bad working practices, induces the premature failure of the components.
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Chapter I - Introduction
Besides the direct costs of the molds, there is a direct impact on the production due to the
increasing machine down times, higher process instability, poor product quality and higher
maintenance costs. Therefore, any optimized solution to achieve the best performance of
molds should be focused on the proper selection of the base material, the coating and its
deposition process. Unfortunately, the lack of credible solutions have limited the application
of coated components in glass industry, hiding the great relevance and impact they could have
for the industrial development in this area. Thus, this thesis is devoted to the improvement of
surface properties and integrity of molds, in order to improve their durability, through either
the optimization of conventional coatings currently applied on mold surfaces or the
development of new coating systems deposited by emergent technologies. Both cases show
potential to over-support the in service conditions of the molds, allowing improving their
ratability, lifetime and performance.
Presently, the molds for glass industry are manufactured either in low cost materials,
such as cast iron and low carbon steel, offering a good relative performance at high
temperature, or in copper-nickel alloys (bronze), due to their excellent thermal conductivity
combined with reasonable hardness and wear resistance, which allow an increasing
production rate of glass containers. In some cases, specific zones of the molds are coated,
mainly with Fe, Co or Ni-based alloys, with the aim of increasing their service life, owing to
their excellent oxidation and mechanical performance under high temperature conditions.
Several hardfacing and thermal spraying processes have been used to coat the molds: i)
Plasma Transferred Arc (PTA), and Gas Tungsten Arc Welding (GTAW) and ii) Flame
Spraying (FS) and High Velocity oxy Fuel (HVOF), respectively. Hardfacing processes have
been preferentially used in the protection of molds surface, due to the strong metallurgical
bond promoted between the coating and substrate, condition required to avoid the catastrophic
failure in service. In thermal spraying processes only a mechanical bonding is established.
Among the hardfacing processes referred above, PTA has been the most versatile and
used, due to its ability to deposit very thick coatings with low porosity and high production
rates. However, the final quality of coatings is largely influenced by the process parameters;
their incorrect selection leads frequently to problems such as the formation of porosity or
cracks in both the coating and the substrate, lack of adhesion of the coating to the substrate
and changes in the chemical composition of the coating metal alloy, compromising the
performance in service. Hence, a reliable understanding of the influence of the deposition
parameters on the coatings properties is required, in order to optimize the lifetime
expectancies and performance of the coatings. Presently, the tendency is to increase the
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Chapter I – Introduction
dilution of the substrate to avoid adhesion problems and, therefore, not to compromise the
molds life, even if the mechanical properties, oxidation resistance and wear performance are
deteriorated. This will be the first issue in discussion in this thesis: how the increase in
dilution of the base material (cast iron), induced by the change in arc current, influences the
properties of Ni-based coatings deposited by the PTA process.
An alternative route to improve the quality and performance of molds is the use of
new, thick or thin coatings deposited by emerging technologies that can lead to significant
improvements on the mechanical properties. Atmospheric Plasma Spraying (APS) and
Physical Vapor Deposition (PVD) are good examples of these deposition processes which
allow to achieve either metallic, ceramic or composite coatings. Such technologies and
materials have shown already to be potential candidates for improving the performance of
mechanical components and parts used in other industries where high oxidation, corrosion
and wear resistances are required. When compared to traditional deposition processes (PTA,
GTAW and FS), their main disadvantage is related to their high cost of investment and
production and, in some cases, lower coatings adhesion. Therefore, any development on this
field for molds protection should be focused on solutions that withstand the severe conditions
to which molds are submitted, extending their service life sufficiently, in order to justify their
higher cost. This is the second part of this thesis: new solutions based on thick and thin
coatings, deposited by APS and PVD processes, respectively. In the first case, the influence of
nanostructured ZrO2 additions to Ni-based alloy coatings deposited by APS process was
studied. The idea of this study was to take advantage of the high hardness and oxidation
resistance of ZrO2 to improve the oxidation resistance and the mechanical and wear properties
of Ni-based coatings. In relation to other thermal spraying processes, such as HVOF, APS
was selected due to its high potential to deposit materials with high melting temperature
(which is the case of ZrO2), although lower particles velocities are achieved; furthermore, in
relation to the PTA process, APS normally gives rise to coatings with much higher hardness.
Apart from the higher cost, the main disadvantage of APS coatings is the need of their
subsequent fusion when applied in parts for the glass industry, in order to obtain compact and
more performing coatings; the application of this operation is not covered in this thesis.
On the other hand, PVD technique has been extensively used to develop new coating
materials, due to its versatility. This process allows "sweep" the chemical composition of a
coating system composed by two or more chemical elements, by changing the deposition
parameters, producing small quantities of material that can be characterized in order to
optimize the best formulation. Moreover, in comparison to thick coatings, PVD allows the
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Chapter I - Introduction
deposition of films with much better properties, particularly the mechanical properties, for
which it is common to have improvements of two to four times. However, despite the
excellent properties of the films and the versatility of the process, its application in
components for the glass industry has been postponed, since the much smaller thickness, in
the micrometer order, has been seen unattractive to the industrial eyes when compared to
thick coatings. However, the experience gained in the application of such PVD films in other
fields of the metalworking industries (for example in machining, stamping, etc), where the
life-time of the coated parts has been easily increased by one to two orders of magnitude,
allows predicting the success of these films in the protection of molds surfaces for the glass
industry. Over the last decades a large number of films (ceramic and composite) have been
developed for high temperature applications, among which the most studied systems were
nitrides, carbides and oxides of Ti, Cr or Al, whose properties (mechanical, oxidation,
corrosion, wear, etc) have been successfully improved by the addition of other chemical
elements. One of the biggest advantages of thin coatings over thick ones is the possibility of
using ceramics as protection material. In fact, ceramic materials cannot be used in the surface
protection of molds for glass parts, due to their low thermal conductivity, which impede the
efficient cooling of the base material, decreasing the productivity of the process. This
difficulty can be overcome if these materials are deposited as thin coatings.
In this thesis, the addition of a specific element (vanadium) to ceramic films (TiSiN),
deposited by DC reactive magnetron sputtering, was studied. TiSiN is one of the best PVD
deposited coatings for high temperature applications and, with V incorporation, it is expected
to improve the mechanical properties and, at the same time, make them self-lubricant at high
temperatures. Currently, in the glass containers production, the walls of the molds are
frequently wet with liquid lubricant to avoid the adhesion of the melted glass to the molds
surface facilitating the removal of glass containers. Frequently, these lubricants (graphite
based lubricants) contain small amounts of sulfur which combined to oxygen forms sulfuric
acid that erodes the mold surfaces creating holes and, therefore, defects that lead to the
premature failure of the mold. Although V additions have been reported to improve the
mechanical properties and the tribological performance of some binary and ternary coating
systems, at high temperature V rapidly diffuses to the surface leading to the loss of its
lubricious effect after a critical short time. Hence, TiSiN system was selected due to its very
high oxidation resistance, similar to the ternary systems CrAlN and TiAlN, where V
incorporation was already studied, but with much better mechanical properties. Moreover, if
deposited with a nanocomposite structure (TiN grains embedded in a Si3N4 matrix) by
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Filipe Daniel Fernandes
Chapter I – Introduction
changing the thickness of the antidiffusion barrier Si3N4 matrix, the control of V diffusion can
be achieved.
Thus, the objectives of this research can be summarized as follows:
i) To analyze the effect of the arc current variation on the substrate dilution and, therefore, on
the chemical composition, microstructure, oxidation and tribological performance of Ni-based
coatings deposited by PTA process over gray cast iron. The influence of an annealing
treatment on the morphology, mechanical properties, and wear performance of coatings will
be also discussed.
ii) To study the influence of the addition of nanostructured ZrO2 to Ni-based coatings,
deposited by APS process, on their microstructure, mechanical properties and tribological
performance. The properties of the composite coatings will be compared with those of the
unmodified Ni and nanostructured ZrO2 coatings.
Taking into account these objectives, besides this introductory part, this thesis is
organized in six chapters. In chapter II, a historical perspective is presented about the main
problems addressed in molds for glass containers production, as well as on the main materials
and protection processes currently and possible use on their fabrication. In chapter III, IV and
V, the main achievements resulting from the experimental work on the deposition and
characterization of the coatings will be analyzed in relation to the three deposition techniques
above presented. Finally, in chapter VI the outputs of this thesis, a comparative synthesis and
discussion of all the results and future research topics will be presented.
The structure of the thesis is based on a short presentation and synthesis of the
research work performed in the scope of each deposition technique, which is further
supported by the papers already published and/or under submission by the author. The
chapters dedicated to the experimental studies are as follows:
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Filipe Daniel Fernandes
Chapter I - Introduction
Chapter III is dedicated to establish the correlations between the process parameters of
Ni-based coatings deposited by PTA and the chemical composition, morphology, structure,
mechanical properties (hardness), oxidation resistance, mechanisms of oxidation, structural
and morphology evolution after annealing and tribological behavior (evaluated by micro-scale
abrasion and pin-on-disc-tests) of the coatings. This chapter is supported by the papers
presented in annexes A to E.
Chapter IV is focused on the study of the influence of nanostructured ZrO2 additions
on the wear resistance of Ni-based alloy coatings deposited by APS. This analysis includes
the characterization of structure, morphology, mechanical properties and tribological behavior
(evaluated in a reciprocating sliding pin-on-disk equipment and micro-scale abrasion) of the
coatings. This chapter is supported by the paper in Annex F.
Chapter V is concentrated on the analysis of TiSiN films deposited by DC reactive
magnetron sputtering with different Si and V contents. The influence of V on the chemical
composition, structure, mechanical properties, residual stresses, oxidation resistance and
mechanisms of oxidation, thermal stability and tribological performance of the coatings is
addressed. This chapter is supported by the papers in Annexes G to I.
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Filipe Daniel Fernandes
List of papers that are the basis of this thesis
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Filipe Daniel Fernandes
List of papers that are the basis of this thesis
Reprints of the papers were made with the written consent of the Publisher and can be found
in annex.
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Chapter II - Historical Perspective
Chapter II
In the current chapter, the base and coating materials, processes of deposition currently used in molds protection
and the main causes for failure problems are reviewed. Moreover, an overview of possible solutions, regarding
processes and materials, for improving the efficiency of molds is presented.
2. Historical perspective
Glass was discovered by the Phoenicians around the year of 3000 BC. The history tells
that glass was discovered when these explorers lit fires near the beaches and observed the
changes occurring to the sand under the action of the intense heat of the fire. Time passed and
Romans began to master the technique of glass production being able to produce some
rudimentary bottles, intended for the storage of a large variety of liquids. Over the following
centuries, glass containers progressed significantly from the design point of view, having been
regarded as a luxury item in the middle age. At that time, only few wealthiest families used
glass containers to the storage of perfumes and other liquids. The industrialization of glass
containers started in England in the year 1600 when the coal was introduced as a solid fuel.
The glass containers were originally produced manually, but due to the increasing market
demand associated to technological development, in the XIX century glass production started
to be mechanically automated. Glass containers were seen as an ecologic and versatile
product, giving rise to their rapid expansion through the market. Since then, to achieve the
required higher production rates to satisfy the market needs, the use of more advanced
materials was always the preferred solution. During the 20th century, materials science and
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Filipe Daniel Fernandes
Chapter II - Historical Perspective
technology knowledge was introduced in the industry and, from the second half of the
century, researchers and engineers poured a part of the materials development efforts on the
modification of materials surface through coatings application. Their superior properties when
compared to bulk materials, allowed extending the performance and lifetime of components
working in severe conditions. Since then, application of coatings has been one of the preferred
ways to modify the surface of materials permitting to extend their use to new potential
applications.
In glass industry, molds and accessories are in direct contact with the melted glass
(temperature about 1050 ºC) being submitted to very severe abrasion, corrosion, wear and
fatigue at high temperature. Molds, mouthpieces and plungers used for glass production are
shown in Figure 2.1. These mechanical components started to be produced in gray cast iron
and low carbon steel due to their relatively good high-temperature performance, allowing
keeping a fairly acceptable thermal conductivity, at an extremely low cost [1, 2]. In fact, the
functionality of these materials was not only to support the severe environment high-
temperature conditions but also to provide an efficient heat transfer which allows to rapidly
cool down the melted glass in order to either decrease the time of containers production or
obtain products without glass distorting under its own weight [2, 3]. Therefore, later on,
copper-alloys (bronze-aluminum) have been introduced as mold material due to their superior
thermal conductivity combined with good properties of hardness and wear resistance [4].
These materials allowed increasing production rates, despite of their much higher cost when
compared to the previous ferrous alloys. Independently of the mold material, some cooling
channels and fins had been performed in some components in order to help the heat
extraction.
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Filipe Daniel Fernandes
Chapter II - Historical Perspective
The higher oxidation and wear resistance required at high temperatures (HT) for
achieving higher production rates, in order to satisfy market demands, led to the application of
new high performing bulk materials, such as Ni cast alloys, that allowed extend service life of
components [5, 6]. However, despite of their excellent performance at high temperature, their
price of acquisition/production was extremely high which have hindered their widespread
application in glass industry. Nevertheless, their attractive properties at high temperatures
pushed to overcome the cost barrier by their application under the form of thick coatings on
components fabricated in cheaper base materials. Currently, the molds are made in a low
performing low cost material (such as the mentioned above, i.e. cast iron, low carbon steel
and copper alloys) but coated in specific zones with hard and more resistant HT alloys,
allowing to support the harsh service conditions [2, 3, 7, 8]. The base material provides the
necessary mechanical strength to resist to the overall applied loads and the thermal fatigue,
whereas the coatings allow extending these properties to an upper end of their performance
capabilities while protecting them against wear, corrosion, oxidation, abrasion, etc [9, 10].
Even if the material withstands high temperature conditions without coating, the life period of
the component can be enhanced. In conclusion, the advantages of coatings application can be
summarized as follows [9, 11, 12]:
Cost of the coating is significantly lower than that of a component integrally fabricated
in HT alloys.
Unique structures and microstructures can be achieved in coatings which are not
possible in bulk materials.
Very high flexibility concerning the chemical composition of the deposited material
allowing the best selection and consequent optimization for a specific
solicitation/application.
With coatings, surface properties can be improved keeping the required mechanical
properties in the bulk of the structural component.
A wide variety of coating materials, such as cobalt, iron and nickel alloys, has been
successfully used to extend the working life of molds [2, 13-15]. The selection of the
appropriate coating composition depends on the environment to which the coating is exposed
and the substrate on which it is applied. The complexity of interactions between environment,
coating and substrate makes the design and selection of coatings very hard and, in general, a
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Chapter II - Historical Perspective
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Filipe Daniel Fernandes
Chapter II - Historical Perspective
requirements. Whenever a failure occurs, the most important factor is not the cost of the
component but the costs due to machine down time, process instability and poor production.
Therefore a proper design and fabrication procedure of the molds, as well as their appropriate
use in service, should be carefully analyzed. The factors that lead to the molds failure are the
thermal and mechanical strains, the wear plastic deformation, the thermal fatigue and the
action of the oxidative, abrasive and corrosive environments occurring in service conditions
[37-39]. The synergetic effect of the harsh conditions with a faulty design, a defective
material or bad working practices is the perfect combination to induce the premature failure of
components, expressed as: worn-out geometries, plastic deformation, surface and bulk
material cracks, breakdown of the coatings. For example, the presence of sharp corners,
notches and sudden changes in the cross section should be avoided in molds design. They act
as hot spots which potentiate the failure by thermal fatigue, wear and cracking due to the high
temperature there reached [38, 40, 41]. With the components of molds being constantly
heated and cooled down during glass containers production, large thermal gradients are
created causing the mold tension during heating and compression during cooling down [37,
42]. Thus, surface cracks are frequently produced resulting in a poor product surface
finishing. Moreover, if the thermal expansion coefficients of the base materials and the
coatings are very different, cracks can appear on either base material, base material/coating
interface or, even, only in the coating, leading also to the mold failure. Thermal shock is also
reported as a cause for premature mold failure [37, 38]. In conclusion, the increase of molds
lifetime, the improvement in glass containers quality, the decrease in the rejection rate and the
reduction of costs production can only be achieved through a convenient selection of new
materials (substrate and coatings), the optimization of molds geometries, application of
innovative processes of deposition and adoption of new working practices.
As it was referred to in Chapter I, the main processes used to apply coatings to protect
the mold surfaces in glass industry, are: (i) hardfacing processes such as Plasma Transferred
Arc (PTA) or Gas Tungsten Arc Welding (GTAW), and (ii) thermal spraying processes such
as Flame Spraying (FS) and High Velocity Oxy Fuel (HVOF). Ones more than the others, in
general, these processes often lead to problems of adhesion of the coating to the substrate,
high rate of porosity (5-15%), high residual stresses, excess of dilution with the consequent
changes in microstructural and functional properties of the coatings and base materials, in all
cases compromising the performance of the molds and the final quality of the produced parts
[43-47]. FS is the oldest process used for depositing coatings on molds for the glass industry,
since it is simple, easy to be applied and of low cost when compared with other deposition
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Filipe Daniel Fernandes
Chapter II - Historical Perspective
methods. However, the adhesion of the coating to the substrate is very low and it presents a
high level of porosity, which frequently reduces significantly the life time of the molds in
service [9]. To avoid adhesion problems, hardfacing processes have been preferentially used
(GTAW and PTA), since they promote a strong metallurgical bond between the coating and
substrate, condition required to avoid their catastrophic failure in service. This technique
gives rise to much better results than thermal spraying processes, for which the adhesion is
only due to a mechanical bonding. GTAW process deposits thicker coatings with lower
porosity and high production rates than flame spraying [48]. However, it is generally
restricted to flat or horizontal surfaces. On the other hand, PTA process produces better
results than GTAW in terms of efficiency and quality of the coatings, at a lower cost. It is a
very versatile process allowing depositing any kind of material, with lower porosity at a
higher production rate than FS and GTAW [49-52]. This process assures a deep penetration of
the coating material into the bulk substrate with a narrow heat affected zone and a relatively
small weld bead due to its high velocity plasma jet when compared to GTAW. However, it is
not portable and cannot be used to coat complex forms, being its main disadvantages.
Presently, PTA is being increasingly used in components for glass industry, mainly on the
mold edges, zone submitted to a high abrasion by the hot plasticized glass. Furthermore,
despite the good adhesion promoted by the hardfacing processes due to the formation of a
metallurgical bond, in some cases this can be a disadvantage. In fact, to reach high adhesion
values it is frequently necessary to increase the level of dilution of the base material with the
coating. Such a procedure leads to changes in the chemical composition of the deposited filler
metal, which can compromise the properties of the coatings and their performance in service.
Hence, a proper selection of the deposition parameters should be carefully done to accomplish
the commitment between the requirements of mechanical strength, corrosion/oxidation
resistance and adhesion [44]. Normally, coatings with thickness in the range of 2 - 5 mm are
applied in the molds surfaces with these processes.
HVOF process provides coatings with higher hardness than PTA do, since it allows
higher cooling rates which promote refined microstructures [9, 53, 54]. Furthermore, owing to
the high energy of the particles jet, better mechanical interlocking can be achieved with the
consequent increase of the bond strength of the coatings to the substrate in relation to FS
process; however, it continues significantly lower than that of the hardfacing processes [55].
As this process is based on the spraying of powders in the melted state over a cooled
substrate, the heat affected zone of the base material remains very low, avoiding the formation
of brittle phases resulting from the metallurgical melting process occurring frequently in
14
Filipe Daniel Fernandes
Chapter II - Historical Perspective
hardfacing processes, which lead very often to failure problems by cracking. The main
disadvantage of this process is the need of fusion of the coating after deposition, since the
rapid cooling down of the powders from the melted state does not allow a good intermixing
among them and a “poor” cohesion of the particles is normally achieved. HVOF allows to
deposit uniform coatings (without dilution) keeping heat extraction conditions uniform in the
entire molds surface, which is a required condition to evenly cool down the glass containers
and, therefore, to produce glass parts without defects. HVOF is being used to coat critical
components in the glass industry, for example the plungers shown on the right in Figure 2.1,
where an efficient and uniform heat transfer is required [2, 3, 8]. The typical thickness of the
coatings for molds surface protection is in the range of 100 - 500 µm. A summary of the main
and typical characteristics of thermal spraying and hardfacing processes is shown in Table
2.1.
Table 2.1 - Comparison of the characteristics of several hardfacing and thermal spraying processes [9, 55-60].
Spray gun Particle Relative Porosity
Dep. Material Temperature reached in Dep. rate
temperature velocity bond level
process feed type the electric arc (0C) (Kg/h)
(0C) (m/s) strength vol.%
FS Powder 3000 - 30-120 Fair 10 - 15 1 - 10
GTAW Wire - 6000-11000 - Excellent 0.01 - 20 6
PTA Powder - 12000-18000 - Excellent 0.1 - 1.5 8
Good to
HVOF Powder 4000 - 2000 0.1 - 2 2 - 12
Excellent
15
Filipe Daniel Fernandes
Chapter II - Historical Perspective
adhesion to the substrate. Despite these disadvantages, APS deposition of cermet and ceramic
coatings was shown in some cases to lead to much better properties than when they are
deposited by HVOF [59]. Thus, the use of the APS process to deposit either the previous
referred cermet coatings (Ni-based alloys reinforced with stable carbides and oxides), which
are normally used to improve the wear behavior, or new cermet coating systems, should be
taken into account. In spite of the extremely high number of studies regarding the effect of
carbide and oxide additions on the tribological behavior of Ni-based alloys, no reports have
been published regarding the effect of nanostructured zirconia additions on their properties.
Presently, pure tetragonal zirconia coatings are one of the main materials used on the turbine
blades surface protection due to their extremely high performance at high temperature [63-
66]. As in glass industry, such material cannot be used as pure coating to protect the mold
surfaces due to its very lower thermal conductivity; however, owing to its high level of
hardness, wear and oxidation resistance, its use as reinforcement material in Ni-based alloy
coatings should be considered.
PVD is a process of deposition of thin films (thickness in the range from few to
thousands nanometers) that produces coatings with outstanding mechanical properties. It is
common in this process, depending on the coating system, to deposit films with two to four
times better mechanical properties than those of the thick coatings described above. The
application of this technique to protect molds for the glass industry is still embryonic, since
their much smaller thickness seems, at industrial eyes, unpromising and unattractive when
compared to that of thick coatings. However, in other industries, where parts are submitted to
extreme high temperature and mechanical loading conditions, these coatings have shown to
successfully increase the life-time and performance of the coated parts, in some cases in more
than one or two orders of magnitude [67-73], allowing predict a potential use of these films to
protect glass molds surfaces. From the development point of view, this technique has been
shown to be very versatile. It easily allows scan the chemical composition of a coatings
system, containing one or more elements, permitting the optimization of the deposition of a
specific film, i.e. the coating with the best compromise between e.g. mechanical, tribological,
oxidation and corrosion properties. Moreover, it allows overcoming the problems of using
ceramic coatings in molds protection. In fact, the lower thermal conductivity of the ceramics
materials can be avoided if the thickness of the non-thermal conductive film is very low. The
main disadvantages of this process are its high investment and production costs, although in
some cases the latter can be decreased. In fact, using PVD process some steps of the molds
production can be removed (for example post-deposition machining tasks, see Table 2.2 for
16
Filipe Daniel Fernandes
Chapter II - Historical Perspective
more detail) and, depending on the mold dimensions, several parts can be coated at the same
time. Therefore, sometimes the final cost of molds production can reach values similar to
those achieved with thick films processes. Finally, it should be taken into account the costs of
glass containers production. On the one hand, the much better properties of thin films in
comparison to thicker ones, may envisage an extended lifetime of the components and
therefore, even if the cost of mold surface protection is higher, the final production costs can
decrease. On the other hand, due to the very low thickness, the possibility of an increased heat
extraction, during glass forming process, gives rise to either important savings in energy,
decrease of the production times or improvement of the final quality of the produced parts.
Coating deposition x x x x x x
Machining x x x x x -
Cleaning x x x x x -
yes if cermets -
Remelting - - - x
no if ceramics -
Polishing - - - x x -
In the last decades, the development of thin coatings for mechanical applications was
concentrated in solutions that could ensure simultaneously the specifications required for a
longer lifetime and an increased productivity, high hardness and high toughness, associated
with high thermal resistance, which should accomplish the desired excellent tribological
behavior at high temperatures. The first generation of hard coatings was concentrated in
single carbides, nitrides and oxides (e.g. TiN, WN, TiC, AlN, WC, Al2O3) [67, 74-76]. The
following generations comprised the mixed compounds, such as TiCN, WCN, (TiAl)N,
TiSiN, etc. [77-81] and the multilayer and multiphase structures (TiN/NbN, TiN/AlN, TiN/W,
WN/Ti, TiSiN, TiBN) [80, 82-84]. Finally, super hard coatings were developed as for
example, c-BN, c-B4C, etc [85, 86]. In all these cases, the extension of the lifetime and
performance of the coated components was successfully achieved.
The development process of hard coatings based on transition metal nitrides started
with single nitrides, such as TiN and CrN, which have been widely used as protective
17
Filipe Daniel Fernandes
Chapter II - Historical Perspective
coatings [87-90]. However, the need for using these coatings in more severe and demanding
environments (particularly, at high temperature applications) led to their alloying with
chemical elements that could in anyway improve their thermal behavior; here multi-element
nitride films appeared [91, 92]. The most known of these is TiAlN, which has higher
mechanical strength and oxidation resistance than TiN [93, 94]. The entire range of Ti/Al
ratios was scanned and, globally, the use of very high Al/Ti ratio gives rise to a significant
improvement of the oxidation resistance at high temperatures [95, 96]. With the same aim,
improvement of the oxidation resistance, the addition of Cr to TiN film has also been
extensively studied [92, 97, 98]. Cr was better in some situations, Al or Si were in others.
Therefore, development of quaternary (e.g. TiAlCrN, TiAlSiN, CrAlSiN) and higher (e.g.
TiAlCrSiN) systems, was performed by the incorporation of other elements such as Si, B, Cr,
Al [71, 99, 100]. Among the ternary coatings, TiSiN was one of the most extensively studied
[101, 102]. Depending on the deposition conditions, these coatings have been reported
consisting of: i) nano-sized TiN crystallites surrounded by a Si-N amorphous matrix [103-
105], or ii) substitutional solid solution of Si in TiN structure (not predicted by the Ti-Si-N
phase diagram) [103, 106, 107]. Their main advantage in relation to previous ternary coatings
is the much higher hardness values that can be achieved, between 40 and 50 GPa, for the same
level of oxidation resistance, particularly if a nanocomposite structure is deposited [95, 103,
108].
Currently, during glass containers production, liquid lubricants are used to avoid the
adhesion between the molds surface and the melted glass. In service, most of the lubricant is
volatilized, due to the high temperatures, promoting the glass adhesion to the surface and,
thus, accelerating the wear of the component and giving rise to poor quality in the produced
parts. Furthermore, these lubricants contain commonly small amounts of sulfur which,
combined to oxygen, forms sulfuric acid that erodes the mold surfaces creating holes and
leading to the premature failure of the mold. Thus, the use on mold surfaces of solid
lubricants, instead of the liquid ones, is of great interest, since their low vapor pressure and,
hence, sublimation does not occur so easily. A wide range of solid lubricant coatings have
been developed in the last decades and successfully applied in tribological components. Solid
lubricant coatings such as WC/C, MoS2, diamond-like carbon (DLC), hex-BN, etc, as well as
their combination in nanocrystalline or multilayer structures, are some examples [70, 109,
110]. However, considerable degradation of the tribological effectiveness of these coatings at
high temperature has been reported due to their low oxidation resistance. To overcome this
shortcoming, new concepts of lubrication have been proposed. Solid lubricants for high
18
Filipe Daniel Fernandes
Chapter II - Historical Perspective
temperature applications can be divided into three categories [111-113]: (i) soft metals (e.g.
Ag, Cu, Au, Pb, and In); (ii) fluorides (e.g. CaF2, BaF2, and CeF3); and, (iii) metal oxides (e.g.
V2O5, Ag2Mo2O7). These coatings were developed by combining the intrinsic properties of
some binary or ternary coatings, that are resistant to oxidation, with specific elements in order
to get the lubricious properties.
All the above three types of solid lubricants materials plastically deform and/or form
low-shear-strength surfaces at elevated temperatures, characteristics responsible for the
lubricious effect. Among them, special attention has been given to the coatings based on the
formation of lubricious oxides, being vanadium the preferred element, reason why particular
attention has been given to the vanadium-containing coatings with the expectation of
formation of Magnéli phases VnO3n−1. These coatings showed interesting tribological
properties in the temperature range 500 – 700 ºC [70, 114-119]. Various series of V rich
coatings have been developed, such as ternary (V,Ti)N [120], CrVN multilayered AlN/VN
[121] and quaternary single layer or multilayered CrAlVN [122, 123] and TiAlVN [116, 124-
126]. A detailed description of the effect of V additions on the mechanical, tribological and
oxidation, properties of these coatings can be found in a recent paper review published by
Franz et. al [127]. In summary, the most important conclusions of these studies are: (i) the
relative amounts of V2O5 detected at the oxidized surface of V rich films successfully
decreases the wear rate and friction coefficient of the coatings; (ii) a strong out-diffusion of V
significantly degrades the oxidation resistance of the coatings; (iii) the out-diffusion only
permits a short time efficiency of solid lubrication with the V-oxides.
The control of the V out-diffusion is now one of the major challenges to reach an
adequate oxidation resistance and suitable tribological properties without compromising the
original properties of the host binary and ternary films. One of the possible proposals to solve
this problem will be to use coatings systems composed by a dual phase, in which one of the
phases act as a diffusion barrier to vanadium. This can be the case of TiSiN system. If
deposited as a nanocomposite structure (nano-sized TiN crystallites surrounded by an
amorphous matrix of Si3N4), the Si-N phase may work as an anti-diffusion barrier [128] and,
therefore, if V is incorporated in solid solution in the lattice of the TiN grains, the controlled
release of V can be achieved. By tailoring the nanostructure of the Ti-Si-N films, i.e
producing coatings with different Si-N layer thickness, the V out-diffusion can be controlled
as well as its release as oxide in the coatings surface. An alternative method would be to
deposit a multilayer structure with Si-N layers of increasing thicknesses from the top to the
bottom of the coating.
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
Chapter III
3. Effect of the arc current variation on the properties of
Ni-based coatings deposited by PTA process
3.1. Introduction
This chapter is dedicated to the analysis of the effect of the arc current variation on the
microstructure, hardness, oxidation performance, structural and morphology evolution after
annealing and tribological behavior of Ni-based coatings deposited by PTA process, and is
supported by the publications presented in the Annexes A to E.
The effect of the PTA current variation on the microstructure and hardness of the
coatings is presented in annex A. The microstructural changes in the coatings as well as in the
respective heat affected zones were analyzed in this paper. Further, the effect of a post-weld
heat treatment (PWHT) on the microstructure and hardness of the coatings and base material
was also examined as well as their wear performance evaluated in micro-scale abrasion. This
method was selected to reproduce the main wear mechanism (3 body abrasion) identified in
the surface of molds which have been in service for long time, although it does not allow
consider the high temperature effect. An exhaustive investigation of the test conditions that
produce such kind of wear mechanism was carried out and it is detailed in the paper of Annex
D. The wear resistance of the coatings, in particular their abrasion resistance was evaluated
before and after annealing in order to understand the effect of temperature on the wear
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
performance of the coatings (publication in Annex C). The influence of the heat treatment on
the hardness and microstructure of the coatings was also discussed. The high temperature
tribological behavior of the coatings was assessed in a pin-on-disc tribometer (publication in
Annex E). Besides the effect of the increasing dilution on the wear behavior of the coatings at
room and increasing temperatures, the base material used in the molds production (gray cast
iron) was also investigated. The publication in Annex B is dedicated to the study of the effect
of substrate dilution on the oxidation behavior of coatings. The oxidation mechanisms
occurring during coatings heat treatment were also discussed in this paper.
The experimental procedure and conditions adopted for the coatings deposition and
characterization are described in detail for each specific investigation on the papers of the
annexes. In the following sections, the main results achieved are summarized.
Ni-based coatings were deposited by PTA process, using the same process parameters
with exception of the arc current, on flat surfaces of gray cast iron blocks in order to study the
effect of substrate dilution on their properties. The following arc currents were used: 100 A
for specimen 1, 128 A for specimen 2 and 140 A for specimen 4. Further, the edges of a block
with a machined U-shaped groove were coated with the same parameters of specimen 2
(specimen 3).
The effect of the arc current variation on the dilution of the coatings is shown on the
macrographs displayed in Figure 3.1. With the arc current increase in the range 100 - 140 A,
the dilution steadily increased from 28% to 59%. Both flat and U-shaped groove machined
specimens revealed to have similar dilution and macrostructures suggesting that the
deposition conditions on flat surfaces blocks could be extrapolated to molds. As it would be
expected, the increasing dilution had a significant influence on the chemical composition of
the coatings, being the main changes related to the decrease of the Ni content and the increase
of Fe, Si, C and Mo contents, in good agreement with the chemical composition of the base
material.
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Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
Figure 3.1 - Dilution induced by PTA process using different weldments: specimen 1 - 100 A, specimen 2 and 3
- 128A, specimen 4 – 140 A.
After deposition, four distinct regions could be observed on the cross section of the
coatings, as follows: the fusion zone (FZ), which corresponds to the coating material (the
nickel alloy), the partially melted zone (PMZ), which is the area close to the FZ where
liquation can occur during weld, the heat affected zone (HAZ), which is not melted but
undergoes microstructural changes, and the base material (BM) which structure remains
unaffected by the deposition process. In all the cases the fusion zone showed a dense
microstructure, free of microcracks and few solidification voids. Increasing substrate dilution
was shown to reduce the content of porosity in the coatings, correlated with the refinement of
the dendritic microstructure (Figure 3.2). The microstructure consisted of dendrites of a Ni-Fe
solid solution phase aligned along the direction of the heat flow. Furthermore, torturous grain
boundaries were displayed, with C flakes (dark-floret structures) evenly distributed in the
matrix. The large number of C flakes on the microstructure was attributed to the dilution of
the base material. Moreover, the dendritic structure became finer as the arc current and,
therefore, dilution increased.
For higher arc currents the heat input during deposition increased, being expected a
coarsened microstructure; however, an opposite behavior was witnessed. According to the
WDS maps (wavelength dispersion spectroscopy) of the most abundant elements of the
coatings, the refinement of the dendritic structure was attributed to the higher number of
precipitates and C-flakes in the grain boundaries, due to substrate dilution. In fact, these
phases normally segregate in the grain boundaries, which increased the rate of heterogeneous
transformation during solidification, hindering the grain growth. A detailed analysis in the
elements distribution with increasing dilution can be confirmed in Annex A.
According to X-ray diffraction patterns of the as-deposited coatings, the major phases
of their structure were primarily (Ni, Fe) face centered cubic solid solution, with the
occurrence of N3Si, Cr5B3 and Fe3Mo3C phases. This phase distribution was in good
agreement with the elemental maps distribution drawn for the coatings. Moreover, according
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
to WDS analysis, other precipitates such as C-B, and Mo-C were also observed in the grain
boundaries, contributing to the microstructure refinement, although they were not detected by
XRD analysis, probably due to their small size and/or content. Close to the interface
coating/base material, a cellular microstructure, finer than the dendritic structure, was
observed for all the specimens (see Figure 3.2 d) for specimen 1). This refinement may be
attributed to the high solidification rates involved owing to the efficient thermal exchange
ensured by the high volume ratio substrate/coating.
C flakes
Grain boundary
a) b)
Precipitates
Interface
Deposit material
Base
Circular island material
c) d)
Figure 3.2 - Optical micrographs of: a) specimen 1, b) specimen 2, c) specimen 4, d) interface of specimen 1.
Despite the previous refinement of the coating microstructure with increasing dilution,
which would suggest an increase of the hardness, an opposite trend was registered, i.e. the
hardness of coatings decreased with increasing arc current, as it is shown in Figure 3.3. This
behavior is undesirable in terms of wear resistance and could only be explained by the
changes in the chemical composition reported above. In all the cases the hardness throughout
the melted material revealed to be approximately constant.
The highest hardness value was observed close to the interface coating/fusion zone,
more specifically on the PMZ (538 HV0.5 for specimen 1), being lower with increasing arc
current, which may be favorable in terms of the toughness of that region. This zone has been
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
found to be formed by hard cementite and martensite phases which are usually responsible for
the brittle behavior of Fe-C alloys. The hardness of the HAZ was significantly higher than the
base material but lower than the PMZ material, as a consequence of the presence of
martensite and the absence of cementite. The proportion of martensite in the microstructure
decreased with increasing distance from the fusion line and, thus, the hardness dropped down
too.
The microstructure of heat-affected zones and the induced mechanical properties were
directly related to the heat input and, therefore, to the current intensity. The high hardness
values found in the PMZ and HAZ made these regions potentially responsible for many of the
mechanical problems occurring in cast iron welds since they were brittle. Thus, a post-weld
heat treatment (PWHT) at 850 ºC for 1h was performed to the specimens in order to reduce
the presence of brittle phases. As can be observed in Figure 3.4 for specimen 4, the heat
treatment successfully reduced the hardness of the PMZ and HAZ zones down to the base
material values without significant changes of the coating hardness. In this case, PMZ, HAZ
and BM zones were formed basically by a perlitic/ferritic structure, justifying the large scatter
of hardness measurements, as shown in Figure 3.4.
550 PMZ
Fusion line
Specimen 1
500 Specimen 2
450 Specimen 3
Hardness HV0.5
Specimen 4
400
350 Melted HAZ
material
300
250
Base material
200
150
100
-4 -2 0 2 4 6 8 10 12
Distance across the interface (mm)
Figure 3.3 - Hardness profile across the interface of PTA specimens.
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
550 Specimen 4
Fusion line
Specimen 4 after PWHT
500
450 PMZ
Hardness HV0.5
400
350
HAZ
300
Melted material Base material
250
200
150
100
-6 -4 -2 0 2 4 6 8 10
Distance across the interface (mm)
Figure 3.4 - Hardness profile across the interface after PWHT in specimen 4.
Two of the previous coatings (specimens 1 and 3) were held at 800 ºC and 900 ºC in
air atmosphere during 2 h in a thermogravimetric equipment (TGA), in order to study the
influence of the substrate dilution on the oxidation resistance of the coatings. The surface, as
well as the cross-section, of the oxidized coatings were observed by scanning electron
microscopy (SEM) and analyzed with both energy dispersive X-ray spectroscopy (EDS) and
XRD diffraction. The results are fully presented in Annex B; specimens 1 and 3 were termed
as coatings C100 and C128, respectively. The thermogravimetric analysis revealed that
increasing the substrate dilution the oxidation resistance of coatings decreased, as shown in
Figure 3.5. This behavior was correlated to the chemical composition changes occurred due to
dilution; in particular due to the high iron content introduced in the coating. As a
consequence, different oxide phases were detected in the two coatings, although in each
coating the phases indexed after thermal exposure at 800 and 900 ºC were similar. According
to XRD and SEM-EDS analyses performed at the oxidized surface of the coatings, the oxide
phase sets detected at 800 ºC and 900 ºC were: (i) B2O3, SiO2, Cr5B3 and OAlB for coatings
produced with lower dilution and (ii) Fe2O3 (hematite), Fe3O4 (magnetite), spinels of
NiFe2O4, and small amounts of NiO, B2O3 and B6O for higher dilution coatings. In the latter,
carbon particles were also detected.
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
1.0 Specimen 1
0.6
900 ºC
0.4
0.2
800 ºC
0.0
0 1200 2400 3600 4800 6000 7200
a) Time (s)
1.0 Specimen 3
900 ºC
Weight gain (mg/cm )
2
0.8
0.6
0.4
0.2
800 ºC
0.0
0 1200 2400 3600 4800 6000 7200
b) Time (s)
Figure 3.5 - Isothermal oxidation curves at 800 and 900 0C of the: a) coating deposited using 100 A arc current,
b) coating deposited using 128 A arc current.
Moreover, the analysis performed on the cross section of coatings showed different
types of oxide layers after isothermal oxidation. In the coating with lower dilution the oxide
scale showed to be mainly formed by a protective thick layer of Si-O in which small amounts
of a phase rich in Ni, Si, Al and B (Si-O, Ni-Si and Al-B-O phases) were uniformly
distributed. On the top a B rich oxide was detected too. On the other hand, a dual layer
structure was formed in the case of the coating with high dilution: an external layer of Fe2O3,
with small features of Fe3O4 and NiFe2O4 spinel rich phases, and an internal Fe3O4 layer with
small amounts of an evenly distributed dark phase rich in silicon, nickel and iron. Similar to
specimen 1, small amounts of B-O were also detected on the top. Furthermore, between the
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
outer and inner layers, a very thin Si-O layer was also identified. Thermogravimetric curves
showed that the coating with lower dilution displayed a parabolic oxidation weight gain as a
function of the time, typical indicating an oxide growth controlled by a protective Si-O layer.
The coating with higher dilution followed the same trend at 800 ºC; however, at 900 0C two
stages in the oxidation curve were detected, the first with a linear increase in mass gain, has
been shown to be a compromise between the outward Fe diffusion through a growing layer of
Fe2O3 and the loss of carbon by decarburization and formation of CO2. The second step
obeyed a parabolic law starting at the moment that the oxide scale, above described, thickened
to a critical value impedding the C liberation, leaving the ion diffusion through the scale the
only mechanism controlling the mass gain. A detailed description of the oxidation resistance
and the diffusion mechanisms occurring with these coatings can be found in Annex B.
In order to use micro-scale abrasion tests to evaluate the effect of substrate dilution on
the abrasive wear resistance of the coatings, an exhaustive investigation of the test conditions
(different abrasive concentrations, loads and sliding distances) that produces 3 body abrasion
wear mechanism was carried out and detailed in Annex D. It was observed that low loads and
higher fraction of abrasive slurry enable 3-body abrasion, whilst, 2-body abrasion becomes
stable at high loads and low volume fraction of abrasive slurry. Moreover, it was perceived
that the specific wear rate essentially depends on the wear rate mechanisms (rolling or
grooving) involved and not on the test conditions employed, since these do not produce
changes in the wear mechanism; i.e. different test conditions which produce similar wear
mechanism can be used to calculate the specific wear rate of coatings. The influence of the
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
45
da
Wear rate (10 mm /N.m)
ys
da
10
40
5
ys
35
3
da
20
30
-4
25
20
15
10
5
0
Spec. 1 Spec. 2 Spec. 4 Spec. 4
annealed during
Figure 3.6 - Specific wear rate of Ni-based coatings in as-deposited and annealed conditions from Annex A and
C, evaluated in a micro-scale abrasion equipment.
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Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
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Filipe Daniel Fernandes
Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process
removal of the tribo-layer and the depletion of graphite from the wear track. This tribological
behavior is presented and discussed in detail in Annex E.
450
Wear rate (10 mm /N.m) RT
400 0
550 C
350 0
700 C
3
300
250
-6
200
150
100
50
0
Specimen 1 Specimen 4 Cast iron
Figure 3.7 - Variation of the wear rate of the as deposited PTA coatings and the gray cast iron with testing
temperature.
In conclusion, the increase in the arc current, and therefore in the dilution of the
substrate, has detrimental effect on the hardness, oxidation resistance and tribological
behavior of the coatings at room temperature. However, beneficial effect on the high
temperature tribological behavior was observed due to the fast formation of oxide layers
which protect the coating surface against wear. Therefore, the dilution degree should be
optimized in order to get the best compromise between the level of oxidation, wear resistance,
mechanical properties and adhesion to make possible to achieve coatings with the best
performance.
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Filipe Daniel Fernandes
Chapter IV – Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings deposited by APS
Chapter IV
4. Influence of nanostructured ZrO2 additions on the wear
resistance of Ni-based alloy coatings deposited by APS
4.1. Introduction
This chapter is dedicated to the study of the effect of nanostructured ZrO2 additions on
the microstructure, micro-hardness and wear performance of Ni-based alloy coatings,
deposited by atmospheric plasma spraying (APS) on a low carbon steel, and is supported by
the paper presented in Annex F.
The coatings with nanostructured ZrO2 additions were produced by either powders of
Ni-based alloy and nanostructured zirconia mixed by mechanical alloying or separate powders
using a dual powder injection system available at the APS equipment. In both cases, two
contents of nanostructured zirconia were added to the Ni-based alloy: 20% and 40% in
volume. All the properties of the coatings were compared, interpreted and discussed in
relation to either unalloyed nickel or nanostructured zirconia coatings and based on the
deposition parameters. The coatings produced with powders prepared by mechanical alloying
using 20% and 40 % of nanostructured ZrO2 were denominated as “Ni+20% MA” and
“Ni+40% MA”, respectively, and the coatings deposited using the dual system of powder
injection with 20% and 40 % of nanostructured ZrO2 were designated as “Ni+20% Dual” and
“Ni+40% Dual”, respectively.
33
Filipe Daniel Fernandes
Chapter IV – Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings deposited by APS
All the coatings showed a typical morphology of plasma sprayed materials, with pores,
lamellae, and partially/un-melted particles (see Figure 4.1). Ni pure coating displayed a
compact and homogeneous microstructure (Figure 4.1 a)), while pure ZrO2 coating showed an
extremely high level of porosity with some cracks, resulting from the tensile residual stresses
generated during cooling down to room temperature (Figure 4.1 b)). Ni+ZrO2 MA coatings
displayed a homogeneous and compact microstructure, with small zirconia particles evenly
distributed in the matrix, whilst Ni+ZrO2 Dual coatings exhibited a porous microstructure, full
of semi-melted Ni powders with large particles of ZrO2 entrapped in their boundaries
suggesting a brittle behavior. Nanostructured zirconia additions progressively increased the
hardness of the coatings. However, coatings deposited using powders prepared by mechanical
alloying displayed much higher hardness than Dual coatings (7.0 - 7.6 against 5.4 - 5.9 GPa,
respectively). This was in good agreement with their lower porosity, higher level of
compactness and finer and more homogeneous distribution of the nanostructured zirconia. In
relation to the structure, the main phases detected corresponded to the ones identified in both
pure Ni-based and ZrO2 coatings, i.e: Ni, Ni-Cr-Fe, Cr23C6, Cr5B3 and ZrO2.
34
Filipe Daniel Fernandes
Chapter IV – Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings deposited by APS
a) b) Cracks
c) d)
e) f)
Figure 4.1 - SEM morphology of the: a) Ni-based alloy coating, b) nanostructured ZrO2 coating, c) and d)
Ni+20% Dual and Ni+40% Dual, respectively, e) and f) Ni+20% MA and Ni+40% MA, respectively.
The influence of the nanostructured zirconia addition on the wear behavior of a Ni-
based coating was studied using reciprocating tribological testing equipment with a ball-on-
plate configuration. A harmonic wave generated by an eccentric and rod mechanism imposed
a stroke length of 2.05 mm at a frequency of 1 Hz. A soda-lime glass sphere of 10 mm in
diameter was used as counterpart, as the most suitable material to study the interaction of
glass with the coatings. All the tests were conducted at room temperature during two hours.
35
Filipe Daniel Fernandes
Chapter IV – Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings deposited by APS
Four different values of normal load were applied to the coated samples and the volume loss,
specific wear rate and friction coefficient were evaluated.
ZrO2 incorporation successfully decreased the specific wear rate and the friction
coefficient of the Ni-based coating, independently of the deposition procedure adopted (see
Figure 4.2). However, this trend is more accentuated in Ni+ZrO2 MA than in Ni+ZrO2 Dual
coatings. Furthermore, increasing ZrO2 content in MA coatings has a decreasing monotonous
effect on the wear rate whereas an inverse trend was observed in Dual coatings for the highest
ZrO2 content. Several factors, such as the hardness, the microstructure and the wear
mechanisms occurring during sliding, were used to interpret those different trends in the
specific wear rates of coatings, as follows:
(i) Ni+ZrO2 MA coatings showed clean wear tracks with longitudinal scratches
identifying grooving abrasion. Thus, the uniform distribution of small ZrO2 particles,
combined with the higher hardness and toughness of these coatings, impeded the liberation of
large wear debris from the coatings, their plastic deformation and adhesion, leading to low
specific wear rate.
(ii) In the case of Dual coatings, the main wear mechanism observed on the worn
surfaces was of adhesion type, as for pure Ni coating, due to the large areas of nickel exposed
to the counterpart. The semi-melted state of Ni powders, the high level of porosity and the
agglomeration of ZrO2 particles located in-between Ni-based lamellae, made the coatings less
tough, inducing an increase of their wear rate.
(iii) The increase of nanostructured zirconia from 20 to 40% in Dual coatings,
promoted higher levels of brittleness in the Ni lamellae boundaries, increasing the volume
loss of the material.
(iv) Pure ZrO2 coating displayed the highest specific wear rate among all the coatings,
due to its brittle behavior.
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Filipe Daniel Fernandes
Chapter IV – Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings deposited by APS
40
35
Ni + 20% ZrO2 MA
Ni + 40% ZrO2 MA
-5
20
1.5
1.0
0.5
0.0
Ni Dual MA ZrO2
Figure 4.2 - Specific wear rate of coatings calculated from the reciprocating tribological tests.
37
Filipe Daniel Fernandes
Chapter IV – Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings deposited by APS
1.2
Ni
Ni + 20% ZrO2 MA
0.8
-2
0.6
0.4
0.2
0.0
0.0 0.5 1.0 1.5 2.0 2.5
Load x Sliding distance (N.m)
Figure 4.3 - Evolution of the wear behavior as a function of the parameter “normal load × sliding distance”.
4.0
Wear rate (10 mm /N.m)
3.0
3
-3
2.0
1.0
0.0
Ni Ni + 20% Ni + 40%
ZrO2 MA ZrO2 MA
Figure 4.4 - Specific wear rate of coatings, calculated from the slope of a straight line fitted to the points plotted
in Figure 4.3.
38
Filipe Daniel Fernandes
Chapter V - Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
Chapter V
5. Study of the influence of V additions on the properties
of TiSi(V)N films deposited by DC reactive magnetron
sputtering
5.1. Introduction
39
Filipe Daniel Fernandes
Chapter V – Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
Three series of TiSiN films with different Si contents, each series with increasing V
content, were deposited in a d.c. reactive magnetron sputtering machine, set with two
rectangular magnetron cathodes working in unbalanced mode. A Ti target, with holes
distributed throughout the preferential erosion zone, and a TiSi2 composite target were used
for the depositions. The different silicon contents were achieved by changing the power
density applied to each target and the V content was varied by changing the number of high
purity rods of Ti and V placed in the holes of the Ti target. Three V contents were added to
the films by using 4, 8 and 12 pellets of V. In all cases the total power applied to the targets
was 1500 W. To serve as reference, a TiN film was deposited from the Ti target using 2000
W of power. Ti and TiN interlayers were deposited as bonding layers, in order to increase the
adhesion of the coatings. The interlayer also contained V when V pellets are placed in the Ti
target. In all the depositions the time was adjusted in order to obtain coatings with
approximately 2.5 µm of total thickness (including interlayers). The total working gas
pressure was kept constant at 0.3 Pa, using approximately 30 and 17 sccm of Ar and N2,
respectively. A summary of the studied coatings and the nomenclature for them adopted (used
in Annex G) is shown in Table 5.1.
0P V 4P V 8P V 12P V
Power (W)
(0 at.%) (~3 at.%) (~7 at.%) (~11 at.%)
S3-0 S3-4 S3-8 S3-12
Ti 1000 TiSi2 500
Increasing Si
40
Filipe Daniel Fernandes
Chapter V - Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
films showed a columnar micro-structure. As Si content was increased, the columns size and
the surface roughness were firstly reduced and, then, increased for the highest silicon content.
V additions did not change the global type of the cross section morphology.
As a general trend, the structure of the films could be assigned to a fcc NaCl-type
crystalline phase as e.g. TiN (ICDD card 87-0628). The influence of increasing Si content on
the TiN system and the effect of V content on the TiSiN coating with the lowest silicon
content (S1-y series) is shown in Figure 5.1. Concerning TiSiN films, a loss of crystallinity
was observed for the coating with the highest Si content, suggesting an amorphization of the
structure. Furthermore, a shift of the diffraction peaks to higher angles was observed with
increasing Si content, suggesting the formation of a substitutional solid solution. In fact, with
Si in solid solution, its smaller atomic radius promoted the contraction of the TiN lattice and
the consequent shift of XRD peaks for higher angles. Residual stresses could not be the cause
for the peaks shift since they were compressive and similar in all the films (approximately 3
GPa). The formation of a substitutional solid solution was explained by the low deposition
temperature, together with a low substrate ion bombardment, condition that do not provide the
necessary mobility of the arriving species for Si-N phase segregation and consequent
formation of a nanocomposite arrangement, i.e. nanocrystalline TiN grains embedded in a
Si3N4 matrix. V incorporation in TiSiN films also made moving the peaks to higher angles,
behavior once again related to a smaller unit cell due to the replacement of Ti by the smaller
V atoms. In these coatings the compressive residual stresses were also in the range of ~3 to ~4
GPa.
Figure 5.1 - XRD patterns of TiSiN coatings with Si content ranging from 0 to 12.6 at.% and S1-y coatings with
V additions.
41
Filipe Daniel Fernandes
Chapter V – Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
The dependence of hardness and Young’s Modulus on the Si and V contents is shown
in Figure 5.2. With Si incorporation, the hardness of the coatings increased up to a maximum
value of 27 GPa (6.7 at.% Si), trend attributed exclusively to a solid solution hardening
mechanism since, as above referred, the level of compressive residual stresses was similar for
all the films. A further increase in the silicon content led to a decrease in the hardness, in good
agreement with the loss of crystallinity of this coating. The Young´s Modulus monotonously
decreased with Si incorporation, which is related to the progressive importance of Si-bonds in
the TiN lattice.
Concerning the effect of V additions, similar evolution was observed for all the Si
compositions (see Figure 5.2 b) for S1-y series). A slight increase in the hardness and
Young´s modulus was observed for lower V contents, followed by a drop for the highest one.
Similar explanation was suggested, the hardness enhancement was attributed to solid solution
hardening since, again, the compressive residual stress values are very similar. The coatings
with the highest Si content displayed the lowest hardness and Young’s modulus among all the
V rich coatings, in good agreement with their low crystallinity.
48 360
340
Young's Modulus (GPa)
44
320
Hardness (GPa)
40
36 300
32 280
28 260
24 240
20 220
16 200
0 2 4 6 8 10 12 14
a) Si content (at.%)
Figure 5.2 - a) Effect of Si content on the hardness and Young’s modulus for the TiN system. b) Effect of V
content on the hardness and Young’s modulus forcoating S1-0.
42
Filipe Daniel Fernandes
Chapter V - Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
S1-y coatings
48 360
44 340
Hardness (GPa)
40
36 300
32 280
28 260
24 240
20 220
16 200
0 2 4 6 8 10 12 14
b) V content (at.%)
In relation to the adhesion of coatings, only two failure modes were observed on the
scratch track of the coatings; the first cracking and the first chipping. Globally all the coatings
showed to be well adherent to the substrate, exhibiting high critical load values. In general V
incorporation did not have significant influence in the critical loads whereas, globally, the
addition of Si gave rise to a small decrease of their values, particularly for Lc1 (see Figure
5.3).
43
Filipe Daniel Fernandes
Chapter V – Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
S1-4
S3-12
S2-0
56
65
S2-8
S3-4
48
S1-8
60 40
S1-0
55 32
S2-12
24
S3-0
S2-4
50
S3-8
16
45 8
40 0
TiN S2-y series
a)
b) c)
Figure 5.3 - a) Adhesion critical loads values for the dissimilar coatings. Coatings were scratched as the normal
force was increased linearly from 5 to 70N. Lc2 critical load for TiN coating was > 70 N. b), c) First coating
cracking and first chiping observed on the scratch track of coating S2-4, respectively.
The effect of increasing Si and V content on the onset point of oxidation of TiSi(V)N
systems were studied in a TGA equipment. The films were annealed from room temperature
up to 1200 ºC at a ramp temperature of 20 ºC/min. Si incorporation in TiN film strongly
increased the onset of oxidation, being S3-0 sample the most oxidation resistant. On the other
hand, V addition to TiSiN decreased the onset point of oxidation down to a value close to 500
ºC, independently of the Si and V contents, value even lower than for TiN film.
In order to better study the oxidation resistance of the coatings, isothermal annealing at
selected temperatures and times were performed. TiN was annealed at 600 ºC during 30 min,
S2-0 was heat treated at 900 ºC during 1h and finally S2-8 was annealed at three different
temperatures: 550 ºC during 1h, 600 ºC during 30 min and 700 ºC during 10 min. The results
44
Filipe Daniel Fernandes
Chapter V - Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
are plotted in Figure 5.4. When time and temperature are considered, TiSiN is far the best
resistant to oxidation. Furthermore, it is also confirmed that V incorporation decreased
significantly the oxidation resistance of TiSiN coatings, even to lower values than TiN.
0.25
Weight gain (mg/cm ) S2-8 700 ºC 10 min
2
0.20
S2-8 600 ºC 30 min
0.15 S2-8 500 ºC 1 h
0.05
TiN 600 ºC 30 min
0.00
0 1200 2400 3600
a) Time (s)
Oxide scale
Oxide scale
TiSiN film
TiSi(V)N film
Interlayers Interlayers
1 µm 1 µm
b) c)
Figure 5.4 - a) Thermogravimetric isothermal analysis of coatings exposed at different temperatures. Cross
section morphology of coating: a) S2-0 oxidized at 900 ºC during 1 h, b) S2-8 oxidized at 600 ºC during 30 min.
The main oxides detected on the surface of the coatings after annealing can be
summarized as follows: i) in TiN and TiSiN films, TiO2 - rutile and anatase, were the main
oxides detected; however, in the last case, SiO2 was also present, localized beneath the TiO2
layer; ii) in TiSi(V)N at 550 and 600 ºC the main oxides were Ti(V)O2 and α-V2O5, the latter
displaying different orientations depending on the isothermal temperature and time; at 700 ºC,
which is a temperature higher than the melting temperature of α-V2O5, β-V2O5 was formed,
resulting from a reduction process of α-V2O5. In situ XRD analysis performed in TiSi(V)N
coatings in the range from 500 – 750 ºC (see Annex H) allowed to demonstrate that the first
45
Filipe Daniel Fernandes
Chapter V – Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
oxide being formed was Ti(V)O2 phase. TiSiN film exhibited a typical parabolic oxidation
weight gain as a function of the time. According to STEM elemental maps analysis a double
oxide layer was formed in this case: on top a TiO 2 layer followed by a SiO2 layer. Thus, the
excellent oxidation resistance of the TiSiN film was a result of the formation of a continuous
protective SiO2 layer in the interface film/oxide scale (for more detail see Annex H) that
hindered the ions diffusion and, thus, protecting the coating from further oxidation. In the case
of the V rich coating, the isothermal curves tested at 550 and 600 ºC showed two steps: at an
early stage, the weight gain increase over time is linear whilst, in the second stage, a parabolic
evolution can be fitted. For higher temperatures (700 ºC) only the parabolic evolution was
observed but with a strong increase in the mass gain due to the melting of V-O phases.
The structural analysis of the oxide scale when the oxidation temperatures are lower
than the melting point of α-V2O5 showed that a two layer-oxide structure was formed, with a
thick porous inner layer (SiO2 plus Ti(V)O2) and an outer discontinuous layer of α-V2O5. To
the first linear step in the oxidation curves was suggested the growing of Ti(V)O2, as a
conjugation of a fast rate of mass transport in the oxide with a slower surface step
accompanying the gas dissociation, adsorption and subsequent entry of oxygen into the
Ti(V)O2 lattice. In fact, at the very beginning of the oxidation process, due to the high Ti
content, a TiO2 layer started growing. The presence of V ions and its high solubility with Ti
promoted the formation of Ti(V)O2 solid solution, which comprised vanadium cations with
lower oxidation states, as V3+ ions [119, 129]. The substitution of Ti4+ by V3+ ions in the TiO2
lattice would increase the concentration of interstitial metallic Ti4+ ions and, simultaneously,
would decrease the number of excess electrons. The oxidation mechanism was inverted in
relation to Ti-N, being the outwards diffusion of metallic ions (Ti3+ Tin+) more favorable than
the inwards diffusion of O2- ions. There were several consequences of this inversion: (i) the
oxidation rate increased substantially, (ii) the kinetics of the oxidation was not controlled any
more by a diffusion process but by the rate of dissociation, adsorption and combination at the
surface of O2- with metallic ions, and (iii) the continuous and compact Si-rich layer cannot be
formed. α-V2O5 and β-V2O5 phase showed to be to be formed exclusively on the top of the
inner layer due to the V ions reduction on the surface. The parabolic beahaviour after the first
linear increase in mass was explained by the oxide scale thickness increase, making harder the
ion transport. Then, from a threshold thickness on, the ionic diffusion is the controlling step of
oxidation, over the dissociation mechanism, giving rise to the parabolic behavior
characteristic of ion diffusion controlled processes.
46
Filipe Daniel Fernandes
Chapter V - Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
47
Filipe Daniel Fernandes
Chapter V – Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
When tested against HSS ball, the counterparts were covered with both as-deposited
and oxidized material; however, in TiSiN coating higher amounts of big wear debris were
detected, resulting from the lower fracture toughness of this coating. Therefore, a higher
abrasion of the surface was produced with consequent higher wear rate. For V-containing
coatings, the much higher wear rates and friction values measured with HSS in comparison to
Al2O3 balls were attributed to a less V2O5 oxide formation on the worn surfaces.
A detailed characterization of the effect of V additions on the coatings wear and their
correlation to the main wear mechanisms is shown in Annex I.
14 Wear track
Al2O3 ball
Wear rate (10 mm /N.m)
12
HSS ball
10
3
8
-6
0
TiN TiVN S2-0 S2-4 S2-8 S3-8
Figure 5.5 - Wear rate of the coatings tested against Al2O3 and HSS balls.
The effect of V incorporation on the abrasion resistance of TiSiN coatings (series S2-y
films) was also assessed using the same test conditions as for the Ni-based coatings evaluated
in Annex A. However, in this case, lower sliding distances were selected due to a different
wear mechanism and a much lower coating thickness. The volume loss of coatings plotted as
a function of the load times sliding distance parameter is presented in Figure 5.6; the
respective results for the specific wear rate are plotted in Figure 5.7. As for the other
tribological tests, the addition of vanadium gave rise to lower specific wear rates. This trend
was more notorious for coatings with low V content. In all cases, the main wear mechanism,
suggested by the analysis of the wear scars produced by the ball was 3-body abrasion, which
reproduced the wear observed on the surface of real molds after a period of service. In
comparison to pin-on-disc tests, the specific wear rate achieved with this abrasion test was
higher about two orders of magnitude.
48
Filipe Daniel Fernandes
Chapter V - Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering
7.0
S2-0
S2-4
6.0
-4 4.0
3.0
2.0
1.0
0.0
0.5 1.0 1.5 2.0 2.5
Load x Sliding distance (N.m)
Figure 5.6 - Evolution of the wear behavior as a function of the parameter “normal load × sliding distance”.
3.0
Wear rate (10 mm /N.m)
2.5
3
2.0
-4
1.5
1.0
0.5
0.0
S2-0 S2-4 S2-8 S2-12
Figure 5.7 - Specific wear rate of coatings, calculated from the slope of a straight line fitted to the points plotted
in Figure 5.6.
In summary, although the nanocomposite structure was not formed in TiSi(V)N films,
V additions significantly improve their mechanical and tribological properties; however, it has
a detrimental effect on their oxidation resistance due to the outwards V ions diffusion to the
surface impeding the formation of a continuous Si-O protective layer. The improvement of
the tribological properties with V incorporations was related to V2O5 formation on the sliding
contact that acted as a lubricious tribo-film, decreasing the friction and protecting the coating
from wear.
49
Filipe Daniel Fernandes
Chapter VI - Conclusions and Future Research
Chapter VI
6. Conclusions and Future Research
The main aim of the research activity carried out in this thesis was the improvement of
the surface properties of molds used in glass containers production, through the application of
several types of coatings. Two methodologies were assessed: i) the optimization of
conventional Ni-based coatings currently applied on mold surfaces; ii) the development of
new coating systems deposited by emergent technologies. In the first part, the effect of base
material (cast iron) dilution, induced by the change in arc current, on the properties of Ni-
based coatings deposited by the PTA process was investigated, while in the second part two
different topics were explored. Firstly, the influence of nanostructured ZrO2 additions on the
properties of Ni-based alloy coatings deposited by APS process were studied while, in the
second topic, the influence of V additions on the properties of TiSi(V)N thin films, containing
different Si contents, deposited by DC reactive magnetron sputtering, was evaluated. The
coatings have markedly different morphological features, physical, mechanical, oxidation and
tribology properties depending on the deposition technology and, thus, their direct
comparison becomes sometimes without sense. Nevertheless, although individual
comparisons were already drawn in each chapter, some common aspects among different
coatings deserve comparison, mainly those directly related with the potential industrial
applicability, as it will be described below.
51
Filipe Daniel Fernandes
Chapter VI - Conclusions and Future Research
The dilution degree of the substrate promoted by changing the arc current affected all
the functional properties of the Ni-based alloy coatings. Their hardness, oxidation resistance
and tribological behavior were strongly influenced by the micro-structure which consisted
typically in Ni-Fe fcc solid solution phase, with intendendritic phases rich in Cr-B, Ni-Si and
Fe-Mo-C. Although, the number of precipitates in the grain boundaries increased with the
dilution, the hardness of the coatings was reduced due to the growing influence of the softer
material of the substrate. As a consequence, also the abrasion wear resistance, calculated from
tests performed at room temperature with silica slurry, decreased with increasing substrate
dilution. The same trend was observed in another tribological test (pin-on-disc), having been
observed specific wear rates varying as much as two up to three orders of magnitude. At high
temperature, dilution showed a beneficial effect on the wear resistance of coatings, behavior
explained by the decrease of the oxidation resistance due to the incorporation of base material
elements. In fact, due to the lower oxidation resistance of the coatings produced with higher
dilution, an easy and fast formation of a high amount of oxide debris on the worn surface
occurred, which by agglomeration gave rise to a protective tribo-layer improving the wear
resistance of the coating. The post-deposition heat treatment of the coatings induced a
reduction of the hardness restricted to the partially melted and heat affected zones. In
conclusion, the advantages and drawbacks associated with the degree of dilution indicate that
to produce the best and most performing coating compromises should be found between the
mechanical properties, level of oxidation and wear resistance of coatings.
Nanostructured ZrO2 additions to Ni-based alloy coatings had different impact when
mechanically alloyed mixed powders (MA) or separated powders (Dual) were used. MA
powders deposited coatings displayed a homogeneous and compact microstructure with small
zirconia particles evenly distributed in the matrix; contrasting with a porous microstructure,
full of semi-melted Ni powders with large particles of ZrO2 entrapped in grain boundaries,
responsible for a brittle behavior. In both cases, addition of ZrO2 improved the hardness and
wear behavior. The brittleness of the coatings was determinant on the wear resistance;
therefore, the wear rates increased when going from MA powders, through Dual powders up
to pure ZrO2 coatings. All the APS coatings showed higher hardness values than the PTA
deposited Ni-based ones; however, their micro-abrasion resistance was worse, differing one
order of magnitude, fact explained by the low cohesion between powders. In conclusion, for
any possible application of ZrO2-containing coatings in the protection of glass molds, an
additional melting treatment will be required, in order to improve the cohesion and overcome
the brittleness problems.
52
Filipe Daniel Fernandes
Chapter VI - Conclusions and Future Research
Si and V additions to TiN and TiSiN systems, respectively, did not change
significantly either the global type of columnar morphology or the diffraction peaks indexed
to a crystalline fcc TiN-type phase. A shift of the diffraction peaks to higher angles was
always observed with either Si or V additions, indicating the placement of those elements in
substitutional solid solution in the TiN lattice. Therefore, the required nanocomposite
structure needed for the control of vanadium diffusion was not deposited. The presence of Si
and V incorporation in solid solution in TiN gave rise to a slight increase in the hardness,
although for the high contents an inverse trend was observed. In comparison to previous thick
coatings, much higher hardness values could be achieved with the thin films, as shown in
Figure 6.1, which is an indication of a possible compensation of their lower thickness for
glass molds protection.
27.3
30
28 27 27.6
26 PTA Coatings 23.8
24 23.2 23
22 APS Coatings
Hardness (GPa)
20
18 PVD Coatings
16
14
12 10.6
10
8 7 7.6
6 4.9
4 2.9
1.9
2
0
A
A
im 1
S2 (TiS )
)
S2 (T N)
O
-0 iVN
(T N)
4
N
40 rO i
N
% 2 M
N %Z N
2 M
ec n
Zr
Ti
en
-4 iSi
iV
iV
-1 iSiV
Sp ime
S2 T
iS
O
S2 (T
Zr
ec
Sp
2
-8
20
i+
i+
N
Figure 6.1 - Hardness of coatings studied in the aim of this thesis, deposited by different deposition techniques.
53
Filipe Daniel Fernandes
Chapter VI - Conclusions and Future Research
vanadium at the oxidized surface was promptly oxidized forming a lubricious and protective
oxide, the V2O5. Therefore, despite the observed decrease in oxidation resistance of films
containing vanadium, this should not be regarded as a disadvantage from the tribological
point of view, as the V2O5 phase detected can act as a solid lubricant. Among all the studied
coatings, V-containing coatings displayed the worst oxidation resistance (see comparison in
Figure 6.2). Although TiSiN coating was the most oxidation resistant coating, even in
comparison with the Ni-based coatings (specimens 1 and 3) well known for applications at
high temperature, the addition of vanadium to TiSiN degraded completely its oxidation
behavior. At temperatures higher than the melting point of the V2O5 phase (~675 ºC), the
integrity of the V rich coatings is loss due to rapid oxidation. This can be a drawback for the
application of TiSi(V)N coatings in molds surface protection. Thus, further studies to control
the V diffusion and, consequently, to improve the oxidation resistance of the coatings, are
required. The objective will be the deposition of these films with different structures, i.e. as a
nanocomposite (TiVN grains embedded in a Si3N4 matrix) or as a multilayer structure (TiVN
layers alternating with Si3N4 layers).
Specimen 3 900 ºC 2 h
0.25 S2-8 700 ºC 10 min Specimen 1 900 ºC 2 h
Weight gain (mg/cm )
2
0.20
S2-8 600 ºC 30 min
0.15 S2-8 500 ºC 1 h
0.05
TiN 600 ºC 30 min
0.00
0 1200 2400 3600
Time (s)
Figure 6.2 - Comparison of the oxidation weight gain of Ni-based coatings deposited by PTA and TiSi(V)N ones
deposited by DC reactive magnetron sputtering.
54
Filipe Daniel Fernandes
Chapter VI - Conclusions and Future Research
55
Filipe Daniel Fernandes
Chapter VI - Conclusions and Future Research
the production of the coatings, despite the high production costs of sputtered films, the
reduction of the steps in molds production, such as final machining and polishing (see Table
2.2 in chapter II) and the high number of components that can be coated at the same time,
when compared to the technology used for the production of thicker ones, mitigate the costs
of protection of the molds. With the expectancy of increasing lifetime, due to their superior
mechanical and tribological properties, thin sputtered films should be considered in relation to
thick coatings in the protection of molds for glass industry. Specifically, V rich coatings are
of great relevance and promising for molds protection. Furthermore, its application could be
extended to related areas where high temperature resistance and lubricious properties are
required, such as the protection of machining tools for high-speed cutting and dry machining
processes.
4.0
PTA Coatings
3.5
Wear rate (10 mm /N.m)
APS Coatings
3.0
3
2.0
1.5
1.0
0.5
0.0
A
A
1
)
S2 (TiS )
)
4
iN
N
N
N
M
N
S2 2 M
en
en
en
iV
V
iV
S2 (TiS
i
m
im
iS
i+ ZrO
iS
2
im
O
ci
Zr
(T
(T
ec
ec
-0
e
%
Sp
2
Sp
-4
-8
Sp
20
-1
40
S2
i+
N
Figure 6.3- Specific wear rate of coatings calculated from the micro-scale abrasion tests.
56
Filipe Daniel Fernandes
Chapter VI - Conclusions and Future Research
10
9
3
PTA Coatings
6
PVD Coatings
-6
5
4
3
2
1
0
n
)
)
4
iN
ro
N
N
en
en
Ti
iV
iV
iV
ti
iS
Ti
im
iS
as
iS
iS
im
(T
(T
(T
(T
ec
C
ec
-0
Sp
-4
-8
-8
Sp
S2
S2
S3
S2
Figure 6.4 - Specific wear rate of coatings calculated from the pin-on-disc tests.
275
250
Wear rate (10 mm /N.m)
0
225 Tested at 550 C
200
3
0
Tested at 600 C
175
150
-6
125
100
75
50
25
0
n
)
4
ro
N
en
en
iV
ti
m
as
iS
im
ci
(T
C
ec
e
Sp
-8
Sp
S2
Figure 6.5 - Specific wear rate of coatings calculated from the pin-on-disc tests performed at high temperature.
57
Filipe Daniel Fernandes
References
7. References
[2] M. Sarwar, A.W. Armitage, Tooling requirements for glass container production for the
narrow neck press and blow process, J. Mater. Process. Technol., 139 (2003) 160-163.
[3] R. Penlington, M. Sarwar, D.B. Lewis, Application of advanced coatings to narrow neck
press and blow plungers in the glass container industry, Surf. Coat. Technol., 76-77 (1995)
81-85.
[4] J.F. Dakan, D.G. Schmidt, Bronze alloy for glass container molds, Patent US4732602 A,
1998.
[5] Franklin Bronze and alloy company, Glass Mold Accessory Castings, web reference:
http://www.franklinbronze.com/investment-casting/glass-mold-accessory-castings.htm.
[6] Zhoushan Putuo HongYu Mold Co. Ltd., An analysis of glass mold material and its
development, web reference: http://pthymj.com/en/news_view.asp?id=106.
[7] E. Romero, R.J. DuMola, Alloy coating for aluminum bronze parts such as molds, Patent
US5441554 A, 1995.
[9] T.S. Sidhu, S. Prakash, R.D. Agrawal, Studies on the properties of high-velocity oxy-fuel
thermal spray coatings for higher temperature applications, Materials Science, 41 (2005)
805-823.
[10] M.H. Li, X.F. Sun, J.G. Li, Z.Y. Zhang, T. Jin, H.R. Guan, Z.Q. Hu, Oxidation behavior
of a single-crystal Ni-Base superalloy in Air at 800 and 900 ºC, Oxid. Met., 59 (2003) 591-
605.
[11] C.T. Liu, J. Ma, X.F. Sun, Oxidation behavior of a single-crystal Ni-base superalloy
between 900 and 1000 ºC in air, J. Alloys Compd., 491 (2010) 522-526.
59
Filipe Daniel Fernandes
References
[14] Y. Wu, S. Hong, J. Zhang, Z. He, W. Guo, Q. Wang, G. Li, Microstructure and
cavitation erosion behavior of WC-Co-Cr coating on 1Cr18Ni9Ti stainless steel by HVOF
thermal spraying, International Journal of Refractory Metals and Hard Materials, 32 (2012)
21-26.
[15] R.A. Mahesh, R. Jayaganthan, S. Prakash, Oxidation behavior of HVOF sprayed Ni–5Al
coatings deposited on Ni- and Fe-based superalloys under cyclic condition, Materials Science
and Engineering: A, 475 (2008) 327-335.
[16] J.H. Chang, C.P. Chang, J.M. Chou, R.I. Hsieh, J.L. Lee, Microstructure and bonding
behavior on the interface of an induction-melted Ni-based alloy coating and AISI 4140 steel
substrate, Surf. Coat. Technol., 204 (2010) 3173-3181.
[17] W. Li, Y. Li, C. Sun, Z. Hu, T. Liang, W. Lai, Microstructural characteristics and
degradation mechanism of the NiCrAlY/CrN/DSM11 system during thermal exposure at 1100
ºC, J. Alloys Compd., 506 (2010) 77-84.
[18] C. Guo, J. Zhou, J. Chen, J. Zhao, Y. Yu, H. Zhou, High temperature wear resistance of
laser cladding NiCrBSi and NiCrBSi/WC-Ni composite coatings, Wear, 270 (2011) 492-498.
[20] I. Hemmati, V. Ocelík, J.T.M. De Hosson, Effects of the alloy composition on phase
constitution and properties of laser deposited Ni-Cr-B-Si coatings, Physics Procedia, 41
(2013) 302-311.
[21] M.G. Hobby, G.C. Wood, The role of nickel in the high-temperature oxidation of Fe-Cr-
Ni alloys in oxygen, Oxid. Met., 1 (1969) 23-54.
[22] M.J. Tobar, C. Álvarez, J.M. Amado, G. Rodríguez, A. Yáñez, Morphology and
characterization of laser clad composite NiCrBSi-WC coatings on stainless steel, Surf. Coat.
Technol., 200 (2006) 6313-6317.
60
Filipe Daniel Fernandes
References
[25] T. Liyanage, G. Fisher, A.P. Gerlich, Influence of alloy chemistry on microstructure and
properties in NiCrBSi overlay coatings deposited by plasma transferred arc welding (PTAW),
Surf. Coat. Technol., 205 (2010) 759-765.
[26] M.C. Lin, L.S. Chang, H.C. Lin, C.H. Yang, K.M. Lin, A study of high-speed slurry
erosion of NiCrBSi thermal-sprayed coating, Surf. Coat. Technol., 201 (2006) 3193-3198.
[28] D.W. Yun, S.M. Seo, H.W. Jeong, Y.S. Yoo, The effects of the minor alloying elements
Al, Si and Mn on the cyclic oxidation of Ni-Cr-W-Mo alloys, Corros. Sci., 83 (2014) 176-188.
[29] S. Buytoz, M. Ulutan, S. Islak, B. Kurt, O. Nuri Çelik, Microstructural and wear
characteristics of high velocity oxygen fuel (HVOF) sprayed NiCrBSi-SiC composite coating
on SAE 1030 steel, Arabian Journal for Science and Engineering, 38 (2013) 1481-1491.
[30] C. Guo, J. Chen, J. Zhou, J. Zhao, L. Wang, Y. Yu, H. Zhou, Effects of WC–Ni content
on microstructure and wear resistance of laser cladding Ni-based alloys coating, Surf. Coat.
Technol., 206 (2012) 2064-2071.
[32] Q.Y. Hou, Z. Huang, J.T. Wang, Influence of nano-Al2O3 particles on the microstructure
and wear resistance of the nickel-based alloy coating deposited by plasma transferred arc
overlay welding, Surf. Coat. Technol., 205 (2011) 2806-2812.
[33] H.-y. Wang, D.-w. Zuo, M.-d. Wang, G.-f. Sun, H. Miao, Y.-l. Sun, High temperature
frictional wear behaviors of nano-particle reinforced NiCoCrAlY cladded coatings, Trans.
Nonferrous Met. Soc. China, 21 (2011) 1322-1328.
[35] Ş. Yılmaz, An evaluation of plasma-sprayed coatings based on Al2O3 and Al2O3-13 wt.%
TiO2 with bond coat on pure titanium substrate, Ceram. Int., 35 (2009) 2017-2022.
[36] R. Vaßen, M.O. Jarligo, T. Steinke, D.E. Mack, D. Stöver, Overview on advanced
thermal barrier coatings, Surf. Coat. Technol., 205 (2010) 938-942.
61
Filipe Daniel Fernandes
References
[37] C. Chen, Y. Wang, H. Ou, Y. He, X. Tang, A review on remanufacture of dies and
moulds, Journal of Cleaner Production, 64 (2014) 13-23.
[38] S. Jhavar, C.P. Paul, N.K. Jain, Causes of failure and repairing options for dies and
molds: A review, Engineering Failure Analysis, 34 (2013) 519-535.
[40] G. Pantazopoulos, S. Zormalia, Analysis of the failure mechanism of a gripping tool steel
component operated in an industrial tube draw bench, Engineering Failure Analysis, 18
(2011) 1595-1604.
[41] G.C.R. Moura, M.T.P. Aguilar, A.E.M. Pertence, P.R. Cetlin, The materials and the
design of the die in a critical manufacturing step of an automotive shock absorber cap,
Materials & Design, 28 (2007) 962-968.
[42] D. Mellouli, N. Haddar, A. Köster, H.F. Ayedi, Thermal fatigue failure of brass die-
casting dies, Engineering Failure Analysis, 20 (2012) 137-146.
[43] T.C. Hanson, G.S. Settles, Particle temperature and velocity effects on the porosity and
oxidation of an HVOF corrosion-control coating, Journal of Thermal Spray Technology, 12
(2003) 403-415.
[45] M. Pouranvari, On the weldability of grey cast iron using nickel based filler metal,
Materials & Design, 31 (2010) 3253-3258.
[46] S.W. Banovic, I.N. DuPont, A.R. Marder, Dilution control in gas-tungsten-arc welds
involving superaustenitic stainless steels and nickel-based alloys, Metallurgical and Materials
Transactions B, 32 (2001) 1171-1176.
[47] V. Ramasubbu, G. Chakraborty, S.K. Albert, A.K. Bhaduri, Effect of dilution on GTAW
Colmonoy 6 (AWS NiCr–C) hardface deposit made on 316LN stainless steel, Materials
Science and Technology, 27 (2011) 573-580.
62
Filipe Daniel Fernandes
References
[49] A.J. Ramirez, J.W. Sowards, J.C. Lippold, Improving the ductility-dip cracking
resistance of Ni-base alloys, J. Mater. Process. Technol., 179 (2006) 212-218.
[52] A. Gatto, E. Bassoli, M. Fornari, Plasma Transferred Arc deposition of powdered high
performances alloys: process parameters optimisation as a function of alloy and geometrical
configuration, Surf. Coat. Technol., 187 (2004) 265-271.
[53] H.-J. Kim, S. Grossi, Y.-G. Kweon, Characterization of Fe-Cr-B based coatings
produced by HVOF and PTA processes, Metals and Materials, 5 (1999) 63-72.
[55] J.R. Davis, Handbook of thermal spray technology, first edition, ASM International and
Thermal Spray Society, USA, 2004.
[56] W.L. Daugherty, G.R. Cannell, Analysis of porosity associated with Hanford 3013 outer
container welds, Practical Failure Analysis, 3 (2003) 56-62.
[59] B.G. Mellor, Surface coatings for protection against wear, Woodhead publishing in
Materiais and CRC Press LLC, North America, 2006.
63
Filipe Daniel Fernandes
References
[61] R.S. Lima, B.R. Marple, From APS to HVOF spraying of conventional and
nanostructured titania feedstock powders: a study on the enhancement of the mechanical
properties, Surf. Coat. Technol., 200 (2006) 3428-3437.
[63] Y. Bai, Z.H. Han, H.Q. Li, C. Xu, Y.L. Xu, Z. Wang, C.H. Ding, J.F. Yang, High
performance nanostructured ZrO2 based thermal barrier coatings deposited by high
efficiency supersonic plasma spraying, Applied Surface Science, 257 (2011) 7210-7216.
[65] S.M. Naga, 21 - Ceramic matrix composite thermal barrier coatings for turbine parts, in:
I.M. Low (Ed.) Advances in Ceramic Matrix Composites, Woodhead Publishing, 2014, 524-
536.
[66] Y.H. Sohn, E.Y. Lee, B.A. Nagaraj, R.R. Biederman, R.D. Sisson Jr, Microstructural
characterization of thermal barrier coatings on high pressure turbine blades, Surf. Coat.
Technol., 146-147 (2001) 132-139.
[67] J.L. Mo, M.H. Zhu, Tribological oxidation behaviour of PVD hard coatings, Tribology
International, 42 (2009) 1758-1764.
[68] J.L. Mo, M.H. Zhu, B. Lei, Y.X. Leng, N. Huang, Comparison of tribological
behaviours of AlCrN and TiAlN coatings-Deposited by physical vapor deposition, Wear, 263
(2007) 1423-1429.
64
Filipe Daniel Fernandes
References
[74] H.C. Barshilia, A. Jain, K.S. Rajam, Structure, hardness and thermal stability of
nanolayered TiN/CrN multilayer coatings, Vacuum, 72 (2003) 241-248.
[75] D. Lin, B. Yan, W. Yu, Composition and structure characterization of WNx films
produced by RF reactive sputtering, MRS Online Proceedings Library, 187 (1990) 161.
[76] O. Rist, P.T. Murray, Growth of TiC films by pulsed laser evaporation (PLE) and
characterization by XPS and AES, Fresenius' Journal of Analytical Chemistry, 341 (1991)
360-364.
[77] Z. Qi, P. Sun, Z. Wang, Microstructure and mechanical properties of TiCN coatings
prepared by MTCVD, in: J. Luo, Y. Meng, T. Shao, Q. Zhao (Eds.) Advanced Tribology,
Springer Berlin Heidelberg, 2010, 796-800.
[79] G. Pezzotti, I. Tanaka, Y. Ikuhara, M. Sakai, T. Nishida, Evidences for dilute solid
solutions in the Si3N4-TiN system, Scripta Metallurgica et Materialia, 31 (1994) 403-406.
[83] S.H. Yao, Y.L. Su, W.H. Kao, T.H. Liu, Tribology and oxidation behavior of TiN/AlN
nano-multilayer films, Tribology International, 39 (2006) 332-341.
65
Filipe Daniel Fernandes
References
[85] W.L. Wang, K.J. Liao, S.X. Wang, Y.W. Sun, Microstructure and semiconducting
properties of c-BN films using r.f. plasma CVD thermally assisted by a tungsten filament,
Thin Solid Films, 368 (2000) 283-286.
[86] M.D.A. Rahman, N. Soin, P. Maguire, R.A. D'Sa, S.S. Roy, C.M.O. Mahony, P.
Lemoine, R. McCann, S.K. Mitra, J.A.D. McLaughlin, Structural and surface energy analysis
of nitrogenated ta-C films, Thin Solid Films, 520 (2011) 294-301.
[87] R. Buhl, H.K. Pulker, E. Moll, TiN coatings on steel, Thin Solid Films, 80 (1981) 265-
270.
[89] J. Lin, N. Zhang, W.D. Sproul, J.J. Moore, A comparison of the oxidation behavior of
CrN films deposited using continuous dc, pulsed dc and modulated pulsed power magnetron
sputtering, Surf. Coat. Technol., 206 (2012) 3283-3290.
[90] P.H. Mayrhofer, H. Willmann, C. Mitterer, Oxidation kinetics of sputtered Cr-N hard
coatings, Surf. Coat. Technol., 146-147 (2001) 222-228.
[92] J. Krzanowski, D. Foley, The effect of Cr content on the oxidation behavior of Ti-Cr-N
films, Coatings, 4 (2014) 308-319.
[93] Y.C. Chim, X.Z. Ding, X.T. Zeng, S. Zhang, Oxidation resistance of TiN, CrN, TiAlN
and CrAlN coatings deposited by lateral rotating cathode arc, Thin Solid Films, 517 (2009)
4845-4849.
[94] L.W. Ma, J.M. Cairney, M.J. Hoffman, P.R. Munroe, Deformation and fracture of TiN
and TiAlN coatings on a steel substrate during nanoindentation, Surf. Coat. Technol., 200
(2006) 3518-3526.
[95] J.H. Hsieh, A.L.K. Tan, X.T. Zeng, Oxidation and wear behaviors of Ti-based thin films,
Surf. Coat. Technol., 201 (2006) 4094-4098.
[96] D. McIntyre, J.E. Greene, G. Håkansson, J.E. Sundgren, W.D. Münz, Oxidation of
metastable single‐phase polycrystalline Ti0.5Al0.5N films: Kinetics and mechanisms, Journal
of Applied Physics, 67 (1990) 1542-1553.
66
Filipe Daniel Fernandes
References
[97] D.B. Lee, M.H. Kim, Y.C. Lee, S.C. Kwon, High temperature oxidation of TiCrN
coatings deposited on a steel substrate by ion plating, Surf. Coat. Technol., 141 (2001) 232-
239.
[99] T.D. Nguyen, S.K. Kim, D.B. Lee, High-temperature oxidation of nano-multilayered
TiAlCrSiN thin films in air, Surf. Coat. Technol., 204 (2009) 697-704.
[101] S.H. Kim, J.K. Kim, K.H. Kim, Influence of deposition conditions on the
microstructure and mechanical properties of Ti-Si-N films by DC reactive magnetron
sputtering, Thin Solid Films, 420-421 (2002) 360-365.
[104] J. Musil, Physical and mechanical properties of hard nanocomposite films prepared by
reactive magnetron sputtering, in: A. Cavaleiro, J.M. De Hosson (Eds.) Nanostructured
Coatings, Springer New York, 2006, 407-463.
[105] S. Vepřek, S. Reiprich, A concept for the design of novel superhard coatings, Thin
Solid Films, 268 (1995) 64-71.
67
Filipe Daniel Fernandes
References
[111] J.S. Zabinski, J.H. Sanders, J. Nainaparampil, S.V. Prasad, Lubrication using a
microstructurally engineered oxide: performance and mechanisms, Tribology Letters, 8
(2000) 103-116.
[112] M.S. Bogdanski, H.E. Sliney, C. Dellacorte, The effect of processing and compositional
changes on the tribology of PM212 in air, Lubric. Eng. 48 (1992) 675-683.
[113] S.M. Aouadi, H. Gao, A. Martini, T.W. Scharf, C. Muratore, Lubricious oxide coatings
for extreme temperature applications: A review, Surf. Coat. Technol., In Press, DOI:
10.1016/j.surfcoat.2014.05.064
[114] D.B. Lewis, S. Creasey, Z. Zhou, J.J. Forsyth, A.P. Ehiasarian, P.E. Hovsepian, Q. Luo,
W.M. Rainforth, W.D. Münz, The effect of (Ti+Al):V ratio on the structure and oxidation
behaviour of TiAlN/VN nano-scale multilayer coatings, Surf. Coat. Technol., 177-178 (2004)
252-259.
[117] G. Gassner, P.H. Mayrhofer, K. Kutschej, C. Mitterer, M. Kathrein, A New low friction
concept for high temperatures: lubricious oxide formation on sputtered VN coatings,
Tribology Letters, 17 (2004) 751-756.
[118] P.H. Mayrhofer, P.E. Hovsepian, C. Mitterer, W.D. Münz, Calorimetric evidence for
frictional self-adaptation of TiAlN/VN superlattice coatings, Surf. Coat. Technol., 177-178
(2004) 341-347.
[119] A. Glaser, S. Surnev, F.P. Netzer, N. Fateh, G.A. Fontalvo, C. Mitterer, Oxidation of
vanadium nitride and titanium nitride coatings, Surface Science, 601 (2007) 1153-1159.
[120] J.H. Ouyang, S. Sasaki, Tribo-oxidation of cathodic arc ion-plated (V,Ti)N coatings
sliding against a steel ball under both unlubricated and boundary-lubricated conditions, Surf.
Coat. Technol., 187 (2004) 343-357.
68
Filipe Daniel Fernandes
References
[121] J.-K. Park, Y.-J. Baik, Increase of hardness and oxidation resistance of VN coating by
nanoscale multilayered structurization with AlN, Mater. Lett., 62 (2008) 2528-2530.
[123] Y. Qiu, S. Zhang, J.-W. Lee, B. Li, Y. Wang, D. Zhao, Self-lubricating CrAlN/VN
multilayer coatings at room temperature, Applied Surface Science, 279 (2013) 189-196.
[126] Q. Luo, Temperature dependent friction and wear of magnetron sputtered coating
TiAlN/VN, Wear, 271 (2011) 2058-2066.
[127] R. Franz, C. Mitterer, Vanadium containing self-adaptive low-friction hard coatings for
high-temperature applications: A review, Surf. Coat. Technol., 228 (2013) 1-13.
69
Filipe Daniel Fernandes
Annex A
Annex A
F. Fernandes, B. Lopes, A. Cavaleiro, A. Ramalho, A. Loureiro, Effect of arc
current on microstructure and wear characteristics of a Ni-based coating deposited
by PTA on gray cast iron, Surface and Coatings Technology, 205 (2011) 4094-
4106.
71
Filipe Daniel Fernandes
Surface & Coatings Technology 205 (2011) 4094–4106
a r t i c l e i n f o a b s t r a c t
Article history: The plasma transferred arc (PTA) technique is currently used to coat the edges of moulds for the glass industry
Received 16 December 2010 with nickel-based hardfacing alloys. However the hardness and wear performance of these coatings are
Accepted in revised form 3 March 2011 significantly affected by the procedure adopted during the deposition of coatings. The aim of the present
investigation is to study the effect of arc current on the microstructure, hardness and wear performance of a
Keywords:
nickel-based hardfacing alloy deposited on gray cast iron, currently used in molds for the glass industry.
Plasma transferred arc
Ni based alloy
Microstructure, hardness and wear assessments were used to characterize the coatings. Electron probe micro
Microstructure analysis (EPMA) mapping, scanning electron microscopy/energy dispersive X-ray analysis (SEM/EDAX) and
Wear X-ray diffraction (XRD) were used to characterize the microstructure of the deposits. The effect of post weld
heat treatment (PWHT) on the microstructure and hardness was also studied. The typical microstructure of
the coatings consists of dendrites of Ni–Fe, in the FCC solid solution phase, with interdendritic phases rich in
Cr–B, Ni–Si and Fe–Mo–C. Increasing the arc current reduces the proportion of porosity and hardness of the
coatings and modifies their composition due to the increasing dilution of the cast iron. The partial melted zone
(PMZ) had a typical white cast iron plus martensite microstructure, while the heat affected zone (HAZ) had
only a martensite structure. The wear tests showed decreasing wear resistance with decreasing hardness of
the coatings. PWHT reduces the hardness of the PMZ and HAZ but does not significantly alter the hardness of
the bulk coating.
© 2011 Elsevier B.V. All rights reserved.
1. Introduction type and dilution of the substrate. Balasubramanian et al. [5] suggest
that the dilution of coatings should be controlled through the relevant
Traditionally, molds for the glass industry are manufactured from process parameters (transferred arc current, travel speed, powder feed
gray cast iron or copper alloys due to their superior thermal behavior, rate, torch oscillation frequency and stand-off distance). Takano et al. [6]
relatively low cost and excellent thermal conductivity. This is studied the influence of arc current variation, arc constriction and
associated with high hardness and good wear resistance and leads plasma gas flow on the characteristics of Co coatings. They concluded
to a high production rate of glass parts. However, the edges of the that current intensity is the parameter that most significantly affects the
molds made of these materials are sensitive to wear in heavy duty characteristics of the deposits as higher currents increase the dilution of
cycles; so they need to be coated with materials that are more the base material and decrease the hardness of the coatings. Díaz et al.
resistant to high temperature and wear [1]. [7] mentioned that the dilution also increases with gas plasma flow in
The coating of these edges is frequently done by plasma transferred stellite 6 coatings.
arc (PTA), using nickel-based filler metals [2]. Normally the PTA process Nickel-based hardfacing alloys have become increasingly popular
produces very high quality, thicker deposits, offering optimal protection in recent years owing to their excellent performance in environments
with minimal thermal distortion of the parts, low environmental impact where abrasion, corrosion and elevated temperature are factors. The
and high deposition rates in single layer deposits [3]. It allows perfectly microstructure of Ni-based hardfacing alloy deposits has been studied
controlled deposition of alloys on mechanical parts that are subject to using various alloy compositions and different substrates [8–10].
harsh environments, significantly extending their service life [4]. The However, most of the studies consider the deposition of nickel alloys
composition and properties of the coating are greatly influenced by the on stainless steels; there are no references to characterization of such
deposits applied to gray cast iron.
Gurumoorthy et al. [8], who worked with nickel-base superalloys
⁎ Corresponding author at: CEMUC, Department of Mechanical Engineering, University
of Coimbra, Rua Luís Reis Santos, 3030–788 Coimbra, Portugal. Tel.: +351 239 790 700;
deposited by PTA on stainless steel, reported that the microstructure of
fax: +351 239 790 701. the deposit consists of γ-Ni solid solution dendrites, carbides, and
E-mail address: filipe.fernandes@dem.uc.pt (F. Fernandes). interdendritic eutectics composed of γ-Ni and other phases identified as
0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved.
doi:10.1016/j.surfcoat.2011.03.008
F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106 4095
example aluminum, titanium and niobium are added to strengthen the Specimen Specimens 2 and 3* Specimen
material through the formation of γ′ gamma prime (Ni3(Al,Ti)) and 1 4
boron and zirconium are added to improve the creep strength and Main arc current (A) 100 128 140
ductility [9,10]. Powder feed rate (rpm) 20 20 20
This paper details an investigation of the microstructure, hardness and Travel speed (mm/s) 2 2 2
Powder feed gas flow rate (l/min) 2 2 2
3-body abrasion behavior of a nickel-based hardfacing alloy deposited by
Plasma gas flow rate (l/min) 2.2 2.2 2.2
PTA on gray cast iron with different weld currents. Microstructural Shielding gas flow rate (l/min) 20 20 20
changes in the HAZ as well as the effect of the post weld heat treatment on Torch work distance (mm) 13 13 13
the microstructure and hardness of the coatings are also examined. Oscillation (mm) 4 4 4
Preheat temperature °C 480–500 480–500 480–500
Table 1
Nominal chemical composition (wt.%) of substrate and hardfacing alloy.
Base material C Mn Si P S Cr Ni Mo V Ti Fe
Grey cast iron 3.60 0.60 2.00 b 0.20 b 0.04 b 0.20 b 0.50 0.50 0.10 0.20 Balance
Hardfacing alloy C Cr Si B Fe Al F Co Ni
Ni-alloy 0.14 2.45 2.56 0.86 1.08 1.30 0.01 0.08 Balance
4096 F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106
Fig. 1. (A) microstructure of various regions in gray cast iron weld in as-weld condition, (B) partially melted zone, (C) heat affected zone, (D) transition heat affected zone–base
material, (E) base material.
Fig. 2. Dilution induced by PTA process using different weldments: specimen 1 — 100 A, specimens 2 and 3 — 128A, specimen 4 — 140 A.
F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106 4097
Table 3 both measured from the cross section of the coating using Photoshop
EPMA punctual examination of the deposit and base material of specimen 3. image analysis. These images show that as the current is increased in
Elements (% W) the range 100–140 A, the dilution steadily increases from 28% to 59%.
Specimen 3 in the image illustrates the groove shape usually machined
Al Si Ti Cr Fe Ni Mo
in the mold before coating. Comparing specimens 2 and 3, which were
1 1.00 2.52 0.00 2.36 27.37 56.12 0.23 Coating
both coated using the same current, it is possible to see that the groove
2 0.96 3.06 0.00 2.50 26.48 56.11 0.17
3 0.14 0.79 0.06 4.64 16.46 62.24 0.24 produces only a small change in the dilution and the results obtained
4 0.01 2.17 0.00 0.03 91.73 0.07 0.27 Interface from the flat blocks can be extrapolated to molds.
5 0.02 2.05 0.01 0.03 93.81 0.10 0.32 Base material The dilution of cast iron modifies the composition of the coatings,
as illustrated in Table 3, which shows the chemical composition
measured by EPMA at various points of the coating and the substrate
Fig. 1 B) to E) illustrates the microstructures of the zones marked with of specimen 3. Point 1 is located close to the coating surface, point 2 at
the digits 1 to 4 in Fig. 1 A) in higher magnification. the mid-thickness of the coating, point 3 in the coating close to the
Fig. 1 B) shows the transition zone between the coating and the interface and the other two points are in the bulk cast iron; point 4
partially melted zone. Fig. 1 C) illustrates that the HAZ is basically close to the interface and point 5 distant from it. The amount of iron in
composed of a martensitic structure (needle shape), which is found the coating (N16%) is much higher than that provided by the
mainly in the region closer to the fusion line, however, precipitates and hardfacing alloy (1.08%), producing a consequent reduction in the
graphite flakes are also present. The microstructure of the base material is nickel content. According to Ezugwu et al. [9] the increase in iron
principally ferritic with flakes of graphite and precipitates uniformly content in the coatings tends to decrease their oxidation resistance
distributed in the matrix (Fig. 1 E)). Fig. 1 D) illustrates the transition because of the formation of a less adherent oxide scale. The increase in
between the HAZ and the base material. The nature and relative size of Mo content in the coating is also caused by the dilution of the cast iron
those zones are determined by the coating procedure, mainly the heat because it is absent in the hardfacing alloy. Although the true carbon
input during the process, and the composition of the cast iron and content value could be skewed by contamination, carbon is present in
hardfacing alloy used. The constitution of the main zones will be discussed the coating in large quantities, again due to the contribution of the
separately, in order to provide a better understanding of the effect of graphite from the cast iron. Thus, increasing the arc current will
current on their characteristics. increase the dilution and the chemical composition of the coating will
change accordingly.
3.1.1. Melted zone All coatings contained small random pores throughout the cross
Fig. 2 illustrates the effect of the current on the dilution produced section, as shown in Fig. 3. According to F. A. L. Dullien [13], the
in each sample. The dilution is defined as the proportion between the “volumetric” and “area” porosity proportions are equivalent for random-
area of melted base material and the total area of the melted zone, structured porous media, which agrees with our results. The area of the
Fig. 3. Aspect of the coatings of: A) specimen 1, B) specimen 2, C) specimen 4, and D) fusion interface of specimen 1.
4098 F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106
Fig. 4. Optical micrograph of: A) specimen 1, B) specimen 2, C) specimen 4, and D) interface specimen 1.
pores visible under the optical microscope was calculated using the ImageJ the current was increased due to greater heat input. However this was
program (developed at the National Institutes of Health, United States). not observed. Fig. 4 shows that the dendritic structure becomes finer
The proportion of porosity calculated is very low for all specimens; 0.35% as the current increases. For example, the average ternary dendrite
for specimen 1 with an average pore size of 1.45±0.58 μm, 0.32% for spacing for the specimens 1 and 4 is 15.4 and 11.2 μm, respectively. As
specimen 2 with an average pore size of 1.61±0.65 μm and 0.28% for explained below, this refinement can be related to the change in the
specimen 4 with an average pore size of 1.37±0.40 μm. This decrease in composition of the deposited material due to its increasing dilution as
the proportion of micropores in the coatings as current increases can be the arc current increases.
correlated with the refinement of the dendritic microstructure mentioned The phenomenon of circular “islands”, which frequently occur close to
below. The decrease in dendrite size reduces the amount of liquid metal the interface of the coating, is common to all the specimens (see Fig. 4 D)
trapped by dendrites during solidification, reducing the number of voids for specimen 1), These islands are characterized by the total absence of
in the microstructure. Charmeux el al [14] mentioned this mechanism in graphite. According to the Commersald company [15] they are caused
the solidification of Al micro castings. On the other hand the cast iron also when powder grains not melted by the plasma arc subsequently melt in
exhibits significant porosity, with pores reaching 17 μm in diameter, as contact with the pool. Near the interface a cellular microstructure finer
illustrated in Fig. 3 D). than the dendritic structure was observed for all the samples, as illustrated
The typical microstructure of the coatings is shown in Fig. 4. The in Fig. 4 D) for specimen 1, where the grain size is approximately 12 μm.
microstructures consist of dendrites of the Ni–Fe solid solution phase, This refinement may be due to the higher solidification rates involved in
with columnar morphology oriented along the direction of heat flow the boundary because of the efficient thermal exchange ensured by the
and torturous grain boundaries. The microstructure contains C flakes high volume ratio substrate/coating of the base material, as suggested by
(dark floret-like structures) equally distributed through the micro- Gatto et al. [3].
structure. The chemical composition of the Ni-based powder (low SEM analysis gives a more detailed view of the microstructure of the
carbon content see Table 1) does not explain the large number of coatings. Fig. 5 shows a SEM image of the microstructure of the coating
these flakes which can only be attributed to the dilution of cast iron of specimen 1. The image shows that in addition to the dendrites of the
induced by the PTA process. Furthermore, small interdendritic Ni Fe solid solution phase and graphite flakes the grain boundary is
precipitates can also be detected. These are possibly carbides, as formed of two phases: a light gray phase (like-grain) and a dark gray
explained below. Fig. 4 A–C) shows that an increase in the arc current one. EDAX spectra showed that the light gray phase and the dark gray
and, therefore, an increase in dilution, allows carbides and C-flakes in phase are rich in silicon and chromium, respectively.
the microstructure to be more easily detected and observed. It was A detailed analysis was carried out to study the main constituents of
thought that there could be some coarsening of the microstructure as the coatings using EPMA. WDS (wavelength dispersion spectroscopy)
F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106 4099
Fig. 5. A) Typical SEM micrograph (etched) revealing the grain boundaries of specimen 1; SEM EDAX spectra of: B) dark gray phase, and C) light gray phase.
maps of the most abundant elements of the coating for all the specimens that the light gray (like-grain) phase is rich in silicon and the dark gray
were obtained. Fig. 6 A) represents a SEM image of the microstructure of phase is rich in chromium, which agrees with the EDAX analysis. Carbon
the deposit for the specimen 1. The respective elemental maps of Fe, Ni, concentrates chiefly in the graphite flakes, as shown in Fig. 6 F), though
Si, Cr, Al, B, C and Mo are shown in Fig. 6 B–I), color-coded such that the at some particular points the increment in the C signal was also observed
lowest concentration of the element analyzed is indicated in purple and in the boundary, mainly coinciding with the B signal (see circled zone in
the highest in red. The results of the WDS maps for specimen 1 reveal B and C maps) suggesting that carbides may have precipitated. Although
that the dendrites are rich in iron, silicon and aluminum, while not visible in the C-map, very small agglomerates of Mo can be detected
boundaries are rich in boron, silicon and chromium. Nickel is distributed in Fig. 6 also suggesting formation of carbides. However, as shown below
almost uniformly throughout the structure, except in areas where by XRD analysis, these carbides were not identified, probably due to
graphite flakes are present. In the boundary chromium appears, to their small size and/or content.
exclude siliconas these elements do not overlap. However, chromium The increase in the arc current and, consequently, in the dilution of cast
and boron do overlap in the boundaries, as illustrated in Fig. 6 E) and H). iron brought some changes in the distribution of the main elements in the
Thus, it can be concluded that the borders of the boundary are rich in Si microstructure, as shown by the WDS maps for specimen 4, illustrated in
whereas the middles of the boundaries are mainly composed of B and Cr. Fig. 7. For this specimen, in comparison to sample 1, the main difference is
Comparing the information from Figs. 5 and 6 it is possible to conclude related to the Ni and Fe signals. It is clear from Fig. 7 that the zones in the
4100 F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106
Fig. 6. WDS maps of the coating of specimen 1: (A) SEM image of the microstructure of the deposit, (B) iron, (C) nickel, (D) silicon, (E) chromium, (F) carbon, (G) aluminum, (H)
boron, and (I) molybdenum.
interdendritic boundaries richer in iron are depleted in nickel and silicon dilution, since they are present only in the cast iron. Collins and Lippold
but that in the boundary these zones overlap with Cr and B. The grouping [16] and Ramirez et al. [17] mention that the presence of precipitates in
of these maps suggests the presence of a Fe–Cr–B phase in the boundaries, the interdendritic regions results in the formation of very tortuous grain
however, as can be seen later from the XRD analysis this is not identified, boundaries and the excess of theses precipitates stops the migration of the
as only a preferential association with the B–Cr phase is indicated, leading grain boundaries. This can justify the lower grain size of the coating
free iron to combine with other elements. The Si rich zones, see Fig. 7 D), deposited with 140 A in spite of the higher energy supplied to the coating
are depleted of Fe and rich in Ni. In specimen 4, as well as the strong signal process. The literature also indicates that an increase in the molybdenum
given by the C flakes tiny molybdenum and carbon traces could be and carbon content improves the rate of heterogeneous transformation
detected in the surrounding material (where C is hardly present), which is during solidification. These elements normally segregate in the grain
a clear contrast with the results from specimen 1. These elements are boundaries, thereby hindering grain growth [18]. Nevertheless, the
distributed preferentially in the boundaries overlapping the Fe map, influence of the much higher number and content of graphite flakes
suggesting the formation of a mixed carbide containing Fe and Mo. The presented in the 140 A sample, these being segregated from the grain
increase in the content of these elements is caused by the increased boundaries during material solidification, should therefore also be
F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106 4101
Fig. 7. WDS maps of the coating of specimen 4: (A) SEM image of the microstructure of the deposit, (B) iron, (C) nickel, (D) silicon, (E) chromium, (F) carbon, (G) boron, and (H)
molybdenum.
considered as an explanation for the lower grain size in this sample. As can the lower intensity experimental XRD peaks. The other peaks can be
be seen in Fig. 4, C flakes are common in the grain boundaries. adjusted either to a Cr-boride phase Cr5B3 (ICDD card 32–0278) or to
The distribution of the chemical elements in the microstructure in the Fe, Mo mixed carbide of M6C (Fe3Mo3C — ICDD 47–1191) type.
specimens 2 and 3 follows the trend mentioned above. The increasing content of this latter phase through samples 1 to 4 sits
Fig. 8 shows the XRD spectra of specimens 1, 2 and 4. The analysis well with the WDS maps presented and interpreted above. Further-
of this data revealed that the major phase present in the coatings is a more, the indexation is in agreement with the results from the
(Ni, Fe) solid solution face centered cubic (ICDD card 47–1405-(111), literature where similar phases were also detected in a nickel alloy
(200), (220), (311), and (222) peaks at 2θ ~ 51.1°, 59.8°, 89.6°, 111.42° deposit on an austenitic stainless steel [8]. In summary, the increase in
and 119.3°, respectively). WDS maps show that in Cr and B rich zones, the arc current increased the dilution of the cast iron, raising the iron,
Si does not exist as it is connected to Ni. The overlapping of Ni and Si carbon and molybdenum content in the coating, promoting dendrite
signals in WDS maps suggests the presence of a Ni–Si phase, such as refinement of the microstructure due to the increase in precipitates
Ni3Si (ICDD card 32–0699) shown in Fig. 8, corresponding to some of (Fe–Mo–C) and C flakes in the grain boundaries.
4102 F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106
Fig. 8. X-ray diffraction pattern of the deposit material obtained for the specimens 1, 2 and 4.
3.1.2. Partially melted zone (PMZ) and heat affected zone (HAZ) susceptible to cracking. The amount of martensite formed depends on
These zones are critical in cast irons since the material can solidify the composition of the cast iron and the thermal history of the zone.
as white iron if cooling is rapid enough. If the amount of graphite
dissolved during welding is high enough it is likely that it will also 3.2. Hardness
give rise to a continuous carbide network. This is undesirable as a
brittle carbide matrix can cause durability problems in the final part. Fig. 11 shows the hardness profiles across the interface for all the
Fig. 9 shows the microstructure of this zone in specimen 3. In this studied specimens. The vertical line in the graph represents the fusion
image, cementite (Fe3C — white phase) can be seen concentrated near line. The hardness through the melted material is approximately
the interface. Fig. 9 B) shows the microstructure of the PMZ zone constant and tends to decrease with increasing weld current. This can
under high magnification. The image displays a large amount of be related to the higher dilution induced by the process.
cementite in the grain boundaries with acicular martensite inside the For all the samples, the highest hardness values were measured in the
grains. Some precipitates, identified as mixed titanium and molybde- area near the fusion line. A maximum hardness of 538 HV was observed
nium carbides (marked respectively with the numbers 1 and 2 in in specimen 1 and decreases with increasing arc current for the other
Fig. 9 and identified in Fig. 10, were also found. As these phases are specimens. As previously documented, the microstructure in that region
hard and brittle this region is sensitive to in-service crack initiation consisted of hard martensite and cementite. The proportion of martensite
and propagation due to thermal and mechanical fatigue, as discussed in the microstructure decreases with increasing distance from the fusion
below. M. Pouranvari [12] also reported that a continuous brittle line and, thus, the hardness decreases too. The microstructure in the HAZ
network of coarse carbides along the weld fusion line may lead to is directly related to the heat input in the process and therefore to the
initiation of cracking. current intensity used.
In the heat affected zone, carbon can diffuse into the austenite The high hardness values found in the PMZ and HAZ make these
during welding and the austenite may subsequently transform into regions potentially responsible for many of the mechanical problems
brittle martensite due to the high cooling rate. Martensite is also experienced in welds in cast iron. The most effective way to reduce
Fig. 9. SEM image of partial melted zone and heat affected zone of specimen 3, (A) transition between coating and the substrate, and (B) magnification of PMZ.
F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106 4103
Fig. 10. Energy dispersive X-ray analysis (EDAX) of the particles shown in Fig. 9 b), (1) titanium carbide with square shape, and (2) molybdenum carbide with rounded shape.
Fig. 12. Hardness profile across the interface after PWHT in specimen 4.
4104 F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106
The wear tests were performed to predict the effect of the different
process parameters on the in-service wear behavior of the coated
components. The ball-cratering test results are shown in Fig. 14, as the
volume of material loss plotted against the product of sliding distance
(l) and normal load (P). A complete reliability analysis is summarized
in Table 4. The wear volume displays linear evolution with P × l, and
the slope corresponds to the specific wear rate. Specimen 1 exhibits
both the lowest wear volumes and the lowest specific wear rate,
therefore, it is the most resistant to wear. This is compatible with the
greatest hardness displayed by the coating of this specimen, as
illustrated in Fig. 11. Samples 2 and 4 display very similar wear
volumes, which is confirmed by similar hardness values; even so,
sample 2 has a specific wear rate slightly higher than sample 4.
Fig. 15 A) shows a SEM micrograph of a wear scar induced by the
ball-cratering device and Fig. 15 B) the surface of a coating of a mold
after a long time in service, which allows pitting to be identified as the
major failure mechanism. The image also shows some cracks that can
be attributed to the localized melting of hard, brittle phases.
Table 4
Results of the linearization analysis.
Specimen 1 6.9 × 10−4 5.2 × 10−5 5.7 × 10−4 to 8.0 × 10−4 0.977
Specimen 2 9.6 × 10−4 6.6 × 10−5 8.2 × 10−4 to 1.1 × 10−3 0.982
Fig. 13. Indentations in base material (specimen 4) after PWHT. A) indentation made in
Specimen 4 8.0 × 10−4 4.6 × 10−5 7.0 × 10−4 to 9.0 × 10−4 0.987
the perlitic structure, and B) indentation made in the ferritic structure.
F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106 4105
Fig. 15. SEM micrograph of: A) wear scars of specimen 1, and B) of the surface of a Fig. 16. SEM micrograph of: A) 3-body abrasion or rolling wear mechanism, and B) 2-
coating after testing in service. body abrasion or grooving wear mechanism.
4. Conclusions References
[10] A.J. Ramirez, J.C. Lippold, Mater. Sci. Eng., A 380 (2004) 245. [15] Commersald, in: C.P. Technology (Ed.) Commersald PTA Technology, Modena
[11] A. Ramalho, Wear 269 (2010) 213. (Italy), http://www.commersald.com, (2009).
[12] M. Pouranvari, Mater. Des. 31 (2010) 3253. [16] M.G. Collins, J.C. Lippold, Welding J. 82 (2003) 288S.
[13] F.A.L. Dullien, Porous media, Fluid transport and pore structure, Academic Press, 1992. [17] A.J. Ramirez, J.W. Sowards, J.C. Lippold, J. Mater. Process. Technol. 179 (2006) 212.
[14] J.-F. Charmeux, R. Minev, S. Dimov, E. Minev, Borovetz, Proceedings of [18] R.M. Imayev, V.M. Imayev, M. Oehring, F. Appel, Intermetallics 15 (2007) 451.
International Conference, 4M2007, Whittles Publishing (2007) 217–220. [19] R.I. Trezona, D.N. Allsopp, I.M. Hutchings, Wear 225–229 (1999) 205.
Annex B
Annex B
F. Fernandes, A. Cavaleiro, A. Loureiro, Oxidation behavior of Ni-based coatings
deposited by PTA on grey cast iron, Surface and Coatings Technology, 207 (2012)
196-203.
87
Filipe Daniel Fernandes
Surface & Coatings Technology 207 (2012) 196–203
a r t i c l e i n f o a b s t r a c t
Article history: The aim of this investigation was to study the effect of PTA current (100 and 128 A) on the oxidation behavior
Received 28 February 2012 of nickel-based hardfacing coatings deposited on gray cast iron. The oxidation behavior of coatings held at
Accepted in revised form 20 June 2012 800 and 900 °C for 2 h in air was studied by thermogravimetry (TGA). The surface, as well as the cross‐section,
Available online 29 June 2012
of the coatings was characterized by scanning electron microscopy combined with energy-dispersive X-ray spec-
troscopy (SEM/EDS) and X-ray diffraction (XRD). TGA results indicate that the coating produced with lower arc
Keywords:
High temperature oxidation
current exhibits more effective oxidation resistance than that produced with higher current. This behavior could
Oxide scales be correlated with the dilution promoted by the PTA process, which changes the chemical composition of coat-
Surface morphology ings. As a consequence, different kinds of oxide scales were detected in each coating after isothermal oxidation. In
Nickel alloys the specimen produced with a lower arc current and lower dilution a protective layer of Si–O is formed, while in
Plasma transferred arc the specimen produced with a higher current two layers could be identified: an external one of Fe2O3, with small
Thermogravimetric measurements features of Fe3O4 and NiFe2O4 spinel rich phases, and an internal one of Fe3O4 with small amount of a dark phase
evenly distributed rich in silicon, nickel and iron. The isothermal oxidation curve at 900 °C of the coating with
higher dilution showed two stages: at an early stage, the weight increase over time is almost linear whereas,
in a second stage, a parabolic law could be fitted to the experimental data. The other specimens followed only
a parabolic law.
© 2012 Elsevier B.V. All rights reserved.
0257-8972/$ – see front matter © 2012 Elsevier B.V. All rights reserved.
doi:10.1016/j.surfcoat.2012.06.070
F. Fernandes et al. / Surface & Coatings Technology 207 (2012) 196–203 197
Table 1
Nominal chemical composition (wt.%) of substrate and nickel alloy.
Base material C Mn Si P S Cr Ni Mo V Ti Fe
Gray cast iron 3.60 0.60 2.00 b0.20 b0.04 b0.20 b0.50 0.50 0.10 0.20 Balance
Hardfacing C Cr Si B Fe Al F Co Ni
Ni-alloy 0.14 2.45 2.56 0.86 1.08 1.30 0.01 0.08 Balance
one is controlled by NiO growth and the second by Al2O3 growth until a examined by scanning electron microscopy with x-ray spectroscopy
continuous Al2O3 layer formed under the previously grown NiO layer. Li (SEM-EDS). Further, the same equipment was used to identify the layers
et al. [12] studied the oxidation of a NiCrAlYSi overlayer with or without forming the oxide scale from the cross‐section of samples. In this case
a diffusion barrier deposited by one-step arc ion plating. They showed the specimens were mounted in epoxy resin, polished and then surfaced
that the duplex coating system exhibits a more effective protection for with a thin layer of gold for perfect observation. X-ray diffraction (XRD)
the substrate, where thin and continuous scales are adhered to the using Co Kα radiation was conducted on the oxidized surface of speci-
overlayer surface, and very limited oxidation and interdiffusion attacks mens. In order to ensure the reproducibility of results, three specimens
are detected. Zhou et al. [13] studied the oxidation behavior of pure and were analyzed for each coating and set of test conditions.
doped nickel alloys with different Co contents in air at 960 °C. They ob-
served that increasing the Co content increases the mass oxidation gain 3. Results
of the coatings.
Since oxidation is one of the main drawbacks that limit the life of 3.1. Microstructure of the as-deposited coatings
glass molds, the aim of this research is to study the effect of PTA current
variation on the oxidation behavior of coatings deposited on gray cast Fig. 1 displays the microstructure of the coatings deposited on the
iron using a nickel-based alloy. The oxidation behavior of the coatings surface of gray cast iron coupons. The coatings show dense micro-
was studied by thermal gravimetric analysis (TGA). The surface and structure without lack of fusion of the substrate, free of microcracks
cross‐section morphologies of coatings after isothermal oxidation were and few solidification voids. Their typical microstructure consists of
observed and characterized by scanning electron microscopy provided dendrites of the Ni-Fe solid solution phase aligned along the direction
with energy dispersion spectrometry (SEM-EDS) and x-ray diffraction. of heat flow. Furthermore, it displays C-flakes (dark-floret like struc-
Further, the high temperature oxidation mechanisms are discussed. tures) randomly distributed in the matrix. The increase in arc current
gives rise to a higher dilution of the gray cast iron, changing the orig-
2. Experimental procedures inal chemical composition of the filler metal. Following the procedure
detailed in reference [9], dilutions of 28% and 54% were measured for
Nickel based Colmonoy 215 (from Colmonoy Company) powder was the coatings deposited using 100 and 128 A arc current, respectively.
deposited by plasma transferred arc (PTA) onto specimen blocks of gray These levels of dilution can explain the increase in C-flakes with in-
cast iron, currently used in the production of molds for the glass indus- creasing arc current, as shown in Fig. 1. Table 3 shows the chemical
try. The deposits were executed using a Commersald Group ROBO 90 composition measured by energy dispersive spectrometry (SEM-EDS)
machine using two different arc currents (100 and 128 amperes). The from an area of 400 × 400 μm, from the cross‐section at the middle
goal of using two different arc currents was to achieve coatings with dis- thickness of each coating. With this procedure it is assured that chem-
similar levels of dilution with the substrate. The nominal compositions ical composition values are measured from a representative volume of
of the cast iron and Colmonoy 215 powder are displayed in Table 1. material in that zone, avoiding erroneous measurements that could
The principal surfacing parameters employed are shown in Table 2. Be- arise from point analyses of the heterogeneous microstructure. As can
fore coating, the blocks of cast iron were induction heated at 480 °C, in be seen, the main differences in chemical composition are related to
order to reduce susceptibility to cracking during coating deposition. the higher iron, carbon and silicon contents achieved with increasing
After coating, specimens containing cast iron and coating were re- dilution. The detailed characterization of the microstructure of each
moved from each block for microstructural analysis. Metallographic coating can be found in a previous publication by the authors [9].
analysis was done using conventional procedures. Small samples for ox-
idation tests were also removed from each coating. The surfaces of these 3.2. Isothermal oxidation in air
samples were ground using number 1000 SiC abrasive paper in order to
produce even surface preparations. Further, the specimens were ultra- The results of the thermo gravimetric analysis performed on the coat-
sonically cleaned in alcohol. ings at different isothermal temperatures (800 and 900 °C) are shown in
Isothermal oxidation tests were conducted at 800 and 900 °C in air Fig. 2a and b for the coatings deposited using 100 and 128 A arc currents,
for 2 h in a thermal gravimetric analysis (TGA) machine. The air flux respectively. From now on and throughout the text, the coatings depos-
used was 50 ml/min and the heating rate up to the isothermal temper- ited using 100 and 128 amperes will be identified as “C100” and “C128”,
ature was 20 °C/min. After thermal exposure the samples were cooled respectively. The C100 coating exhibits parabolic oxidation weight gain
naturally in air to room temperature. The weight gain of samples was as a function of time for both testing temperatures as does the C128 coat-
evaluated at regular 2 s intervals using a microbalance with an accuracy ing at 800 °C. On the other hand, oxidation of the C128 coating at 900 °C
of 0.01 mg. After oxidation the surface morphology of specimens was starts with a linear increase in mass gain but after 1400 s it starts to
Table 2
Main deposition parameters for PTA weld surfacing.
Main arc Powder feed Travel speed Powder feed gas flow Plasma gas flow Shielding gas flow Torch work Oscillation Preheat
current (A) rate (rpm) (mm/s) rate (l/min) rate (l/min) rate (l/min) distance (mm) (mm) temperature,
(°C)
detected at 800 °C and 900 °C were: (i) B2O3, SiO2 and OAlB for coatings
produced with lower dilution and (ii) Fe2O3 (hematite), Fe3O4 (magne-
tite) and spinels of NiFe2O4 for higher dilution coatings. The indexation
of oxide phases of coating C128 is in agreement with the results from
the literature, where similar oxide phases were detected after thermal
exposure [16,17]. Both coatings also display a high intensity peak close
to 60°, which corresponds to the face centered cubic (Ni-Fe) solid solu-
tion (ICDD card 47‐1405). Furthermore, coating C100 displays low in-
tensity XRD peaks which correspond to a Ni–Si phase, such as Ni3Si
(ICDD card 32‐0699), and a Cr-boride phase Cr5B3 (ICDD card 32‐
0278). All these phases are currently identified in this type of micro-
structure [9]. The large amount of iron oxide detected at the oxidized
surface of coating C128 can be attributed to the higher content of iron
in this sample. As referred to before, the higher iron content in this coat-
ing is attributed to higher dilution of the base material induced by the
use of a higher arc current. Comparing the peaks intensity of coating
C128, oxidized at 800 °C and 900 °C, the amount of oxide products
Fe3O4, Fe2O3 and NiFe2O4 increased with temperature. The NiFe2O4 spi-
nel phase (ICDD card 10‐0325), identified at the oxidized surface, results
from the reaction of NiO (which was also identified as an oxide phase on
the surface of coating C128) with the Fe2O3 phase, as reported by Musić
et al. [18]. Furthermore, B-O phases (ICDD card 06‐0297 (B2O3) and
ICDD card 50‐1505 (B6O)) could be identified from the XRD pattern of
coating C128.
Fig. 1. Typical microstructure of: A) coating deposited using 100 A of arc current, B) coating The different chemical compositions of coatings brought about by di-
deposited using 128 A of arc current.
lution led to the formation of different kinds of oxide scales during iso-
thermal oxidation, as shown above. The SEM surface morphologies of
follow a parabolic path. The figures also reveals that for both oxidation the oxidized coatings at 900 °C are shown in Figs. 5 and 6 for coatings
temperatures, coating C100 displays less oxidation weight gain than C100 and C128, respectively. EDS examinations were conducted of the
coating C128. The dissimilar levels of oxidation observed in the speci- oxidized surface of coatings to characterize their composition. Coating
mens can be correlated with the dilution promoted by the PTA process, C100, oxidized at 900 °C, displays two phases: a gray dark phase, identi-
that changed their chemical composition, as shown in Table 3, which fied in Fig. 5a by letter A, and an evenly distributed white phase,
will interfere with the oxidation process. As mentioned before, the indentified by letter B. A magnification of each zone is shown in Fig. 5b
main differences in chemical composition of coatings are related to sili-
con, carbon and, particularly, iron content. Wallwork [14] and Ezugwu
[15] reported that if iron content is increased in nickel-based alloys, it
tends to decrease their oxidation resistance, as is the case in our study.
This behavior was attributed to a progressively higher cation diffusion
rate in the scale. Therefore, increasing the dilution of gray cast iron by in-
creasing the arc current proves to be detrimental, since the oxidation re-
sistance of coatings is reduced due to a further increase of iron content in
the coatings. Moreover, it is observed that the oxidation gain of speci-
mens increases with increasing isothermal temperature. This is an
expected effect since all diffusion phenomena ruling the oxidation pro-
cess are enhanced as the temperature is increased. In conclusion, the
thermo gravimetric results show that increasing dilution decreases the
oxidation resistance of the coatings.
Figs. 3 and 4 show the XRD spectra found at the oxidized surface of
the coatings. The oxide products revealed differences between the two
coatings; however, in each coating the phases produced after oxidation
at 800 and 900 °C are similar, although different oxidation weight gains
were measured. According to XRD patterns, the main oxide phase sets
Table 3
Nominal chemical composition (wt.%) analyzed at middle thickness of coatings depos-
ited using 100 and 128 A arc current.
Elements Ni Fe C Si Al Cr
Coating deposited using 100 A 85.40 8.47 0.30 2.17 0.81 2.86
Fig. 2. Isothermal oxidation curves at 800 and 900° C of: a) coating deposited using 100
Coating deposited using 128 A 51.09 43.29 1.01 2.54 0.50 1.87
A arc current, b) coating deposited using 128 A arc current.
F. Fernandes et al. / Surface & Coatings Technology 207 (2012) 196–203 199
Fig. 3. X-ray diffraction patterns of oxidized surface of coating deposited using 100 A Fig. 4. X-ray diffraction patterns of oxidized surface of coating deposited using 128 A
arc current, after oxidation at 800 and 900° C. arc current, after oxidation at 800 and 900° C.
and c in order to better illustrate the phases. EDS analysis reveals that the could be detected at the oxidized surface of the coating, as shown by
gray phase is rich in boron and oxygen, suggesting that it is a boron peak ID 3. Finally a phase rich in nickel, oxygen, carbon, silicon and
oxide, as identified by XRD analysis. A similar phase was detected by Li boron (see peak ID 4) was detected at the oxidized surface of coating
and Qiu [19], and Lavrenko and Gogotsi [20] after oxidation of boron 128 A explaining the Ni–O, and B–O phases identified in the XRD
carbide powder in air, ranging from 500 to 800 °C and, oxidation of pattern of the specimen.
hot-pressed boron carbide in air up to 1500 °C, respectively. They ob-
served that even if the isothermal oxidantion temperature exceeds the 3.5. Cross‐section of oxidized samples
melting point of the boron trioxide (about 577 °C), after cooling it is pos-
sible detect cristaline boron oxide. According to them, at a temperature The microstructures of cross-sections of the oxide scale of coatings
close to 577 °C the boron oxide melts, covering the oxide surface and C100 and C128, after oxidation and testing at 900 °C for 2 h, are pres-
serving as protective layer. Therefore, oxygen is inhibited from passing ented in Figs. 7a and b, respectively. The oxide scale is about 5 μm
inward and metal atoms from passing outward through the oxide thick for the specimen coated using the lower arc current and 10 μm
layer. Above a threshold temperature, liquid boron oxide starts to evap- thick for that using the higher current. Decarburization of coating
orate, but only after1000 °C does this evaporation becomes significant. C128 close to the surface is also revealed. Comparing this cross‐section
Therefore, when the oxidation temperature is below 1000 °C, during with the oxide scale of the same coating oxidized at 800 °C (Fig. 8) it is
cooling down to room temperature, it is expected that boron oxide possible to conclude that at this temperature there is no decarburization
will solidify on top of the oxidized surface of the material. The white of the specimen. Zones of silicon segregation were found near the inter-
phase appears associated with spalling zones and cracks. The EDS anal- face in both coatings (C100 and C128) annealed at either temperature,
ysis shows that different phases coexist in this zone. A zone rich in as shown, as an example, by the EDS analysis performed on the phase
boron (see peak ID 2) and another rich in silicon (see peak ID 3) were marked with number 1 in Fig. 7c for coating C128 oxidized at 900 °C.
identified, as shown in Fig. 5c. These results matched with XRD analysis The structure of the oxide scale will be discussed separately for each
which suggest that the main phases of the oxidized surface are boron coating, in order to more clearly illustrate its composition and the role of
and silicon oxides, which are responsible for oxidation resistance of thermodynamic stability and transport mechanisms in the establish-
the coating due to their protective character [17,19–21]. Moreover, a ment of layered structures.
floret-like structure appears to be frequently associated with the Fig. 9 shows SEM images in secondary and backscattered electrons of
white phase, as shown in Fig. 5d. The EDS results (see peak ID 4) cross‐sections of the oxide scales of coatings C100 and C128 at 900 °C.
show that the elemental composition of this oxide is rich in aluminium. The oxide layers are complex and composed of a mixture of phases.
Taking into account the XRD results it can be concluded that the main EDS analyses were conducted on the cross‐section of the oxide scales,
phase of the floret-like oxide is OAlB. The morphology and indexation in order to characterize the different oxides formed during isothermal
of this phase named aluminum borate whiskers, matches well with oxidation. Table 4 summarizes the EDS analyses performed at the points
that studied in references [22,23]. marked in the figures. The values plotted in this table indicate the ratio
However, analysis of the oxidized surface of coating C128 at 900 °C between the peak intensity of each element with the peak intensity of
reveals that the oxide morphology is different from that of coating oxygen for each EDS spectrum. This means that at each point analyzed,
C100 (see Fig. 6a). It displays a microstructure with uniformly distribut- a relative increase in this ratio for one particular element is a result of
ed irregular ditches and dissimilar phases. A magnification of the micro- the preferential formation of the oxide of that element over the others.
structure is shown in Fig. 6b to make this clearer and more easily Combining SEM images with EDS results, plotted in Table 4 it can be con-
characterized. The EDS analysis reveals that the phase with irregular cluded that the oxide scale of coating C100 is mainly composed of silicon
ditches (peak ID 1) is rich in iron and oxygen, suggesting an iron oxide (as the phase identified with point 1 suggests), with small
oxide, as identified in the XRD pattern. Further XRD analysis revealed amounts of a dark phase rich in nickel, silicon, aluminum and boron
that the oxidized surface is composed primarily of Fe2O3 with small (see phase identified with point 2) uniformly distributed through the
amounts of Fe3O4. This result is in agreement with that of Guo et al. layer. This agrees well with the XRD pattern and the indexed Si-O,
[16]. It is possible to identify other small features such as another Ni-Si and O-Al-B phases detected. In addition, it is possible to identify a
phase rich in iron (see peak ID 2 relating to the angular particles) for phase rich in boron, normally located at the top of the oxidized surface
which EDS analysis shows the presence of nickel. The combination of (point 3), suggesting a boron oxide, as also revealed by XRD. Moreover,
Ni with the Fe2O3 phase can give rise to NiFe2O4 spinels which were a Ni-rich white phase (see point 4) appears at the interface between
also identified by XRD. Furthermore, particles of carbon (dark phase) the coating and the oxide layer, corresponding to the remaining part of
200 F. Fernandes et al. / Surface & Coatings Technology 207 (2012) 196–203
the original source of elements which diffused outwards to form oxide between the coating and the oxide scale (point 5) is a consequence of the
scales. oxide layer serving as a diffusion barrier for C coming from the decarburi-
The oxide scale of coating C128 is different from that of coating C100, zation of the sample mentioned above. Analogous oxide products were
being composed of two layers. Combining the XRD analysis, with the observed in the cross‐sections of the oxide scale of coating C128 oxidized
SEM-EDS results, the external layer can be identified as Fe2O3, the main at 800 °C.
phase containing small features of Fe3O4 and spinel NiFe2O4. On the A good correlation was found between the scale microstructure and
other hand, the internal layer is composed of a continuous Fe–O rich the isothermal oxidation curve of coating C100 oxidized at 900 °C. The
phase (phase identified with point 2) with small evenly distributed oxidation kinetics of metals and alloys is often controlled by a parabolic
amounts of a dark oxide rich in silicon, nickel and iron, as point 3 of rate law at high temperature [25] due to the presence of protective scales.
EDS analysis shows. The lower iron content in the internal layer in rela- This may well be the case in this specimen, as the evolution of the oxida-
tion to the external one suggests that the internal Fe–O rich phase re- tion weight gain is controlled by the silicon oxide, the main phase present
spects the Fe3O4 oxide. This result is coherent with the oxidation of in the oxide scale, as suggested by Fig. 9. At lower temperatures, a boron
Fe-based materials and so hematite (Fe2O3) and spinel of NiFe2O4 can oxide starts forming on the surface of the specimen due to the high affin-
be expected to form in the region of higher oxygen potential while mag- ity of boron to oxygen. However, for higher temperatures, due to the in-
netite (Fe3O4) will form in the lower oxygen potential region [24]. Simi- tense segregation of silicon through the interface and its great affinity for
lar iron oxide layers were observed by Guo et al., after oxidation of Fe– oxygen, a silicon oxide layer starts to be formed beneath the boron oxide.
36% Ni bicrystals in air at high temperature [16]. Between these outer At the same time small amounts of Ni–Si and O–Al–B phases are formed
and inner layers, a continuous thin layer of a dark phase appears at the due to the diffusion of other elements through the interface. At a temper-
interface, as revealed by SEM. A boron rich phase could be detected at ature of approximately 577 °C, boron oxide melts, covering the oxidized
the top of the outer layer (see point 4), which is in good agreement surface of the coating and inhibiting the transport of species though the
with the boron oxide detected by XRD. Finally, clusters of carbon are dis- melted oxide layer. Above this temperature it is expected that melted
tributed evenly in the oxide layers. The accumulation of C at the interface B2O3 starts to evaporate: however, this vaporization does not
Fig. 5. A) SEM observation of surface morphology of the coating deposited using 100 A arc current, after 2 h oxidation at 900 ° C, B) and C) magnification of the zones A and B iden-
tified in A), D) magnification of the zone C identified in C). SEM/EDS spectra of: E) zone 1, F) point 2, G) point 3, H) zone 4.
F. Fernandes et al. / Surface & Coatings Technology 207 (2012) 196–203 201
Fig. 6. A) SEM observation of surface morphology of the coating deposited using 128 A arc current, after 2 h oxidation at 900°C, B) magnification of the microstructure in A). SEM EDS
spectra of: D) point 1, D) point 2, E) point 3, F) point 4.
Fig. 7. Back-scattered SEM image of the cross‐section of the oxide scale obtained at 900° C of: A) coating deposited using 100 A arc current, B) coating deposited using 128 A arc
current. C) Segregation of silicon in the coating deposited using higher arc current, oxidized at 900° C. D) SEM EDS spectra of point 1.
202 F. Fernandes et al. / Surface & Coatings Technology 207 (2012) 196–203
Table 4
Ratio between the peak intensity of each element with the peak intensity of oxygen (O Ka).
Point B Ka C Ka Fe La Ni La Al Ka Si Ka
significantly influence the slope of the kinetics curve at 800 and 900 °C,
since it only has an important impact at temperatures above 1000 °C
[19,20]. Finally, due to the continuous segregation of silicon, its oxide
layer will progressively thicken. A protective layer was therefore
established and the reaction rate becomes governed by the rate of species
diffusion through this thicker silicon oxide layer, as observed in the cross‐
section of the oxide scale of the specimen. As with coating C100, a strong
correlation can be found between the microstructure of the oxide scale
and the isothermal oxidation curve of coating C128. At 800 °C, parabolic
behavior is displayed; however, at 900 °C, two steps can be detected, see
Fig. 2b. The isothermal oxidation curve at first exhibits a linear increase in
mass gain and then starts to conform to a parabolic law. The microstruc-
tural analysis of the cross-section of the oxide scale revealed that at
800 °C there is no decarburization of the coating, which is not what
was observed at 900 °C. Like the specimen coated using the lower arc
current, at low temperatures boron oxide starts forming on the surface
of specimen due to the high affinity of boron to oxygen. Moreover, for
high temperatures, due to the high iron content, a consistent layer of
Fig. 8. SEM images of the cross‐section of the oxide scale obtained at 800° C of the
Fe2O3 with small features of Fe3O4 and NiFe2O4 is formed, which is thick-
coating deposited using 128 A arc current: A) SEM secondary electron (SE) image, ened by outward Fe diffusion. At the same time the specimen loses car-
B) back-scattered SEM image. bon by decarburization, due to the formation of CO2. However, carbon
Fig. 9. SEM images of the cross‐section of the oxide scale of the coatings deposited using 100 and 128 A arc current oxidized at 900°C: A) and C) in back-scattered SEM images, B) and D) in SEM
secondary electron (SE) images.
F. Fernandes et al. / Surface & Coatings Technology 207 (2012) 196–203 203
has increasing difficulties diffusing outwards through the oxide layer and an internal layer of Fe3O4 film with small evenly distributed amounts
stays either inside or beneath the oxide scale, as shown by morphological of a dark phase rich in silicon, nickel and iron. In summary, an increase
analysis of the cross-section. Therefore, after the first step where the in arc current from 100 to 128 A in the deposition of nickel-rich layers
weight gain is partially compensated by C liberation, the oxidation on substrates of cast iron by the PTA process is very harmful, as it
curve will deviate to a normal parabolic trend. According to the reduces the oxidation resistance of the coatings at high temperature.
Richardson-Ellingham diagram, Si has higher affinity for O than Ni and
Fe, and Fe forms Fe3O4 more easily than Ni forms NiO. Therefore, during
Acknowledgements
oxidation it would be expected that Si-O would be the first layer to be
formed. However, its much lower content in the specimen relative to
The authors wish to express their sincere thanks to the Portuguese
iron, indicates that considerable time is required to form a continuous
Foundation for the Science and Technology (FCT), through COMPETE
and protective SiO2 layer, as suggested by Douglass and Armijo [26]. So,
program from QREN and to FEDER, for financial support in the aim of
after a critical time a layer of SiO2 starts being formed below the Fe2O3
the project number “13545”, as well as for the grant SFRH/BD/68740/
scale, but it never becomes thick enough to determine the oxidation ki-
2010.
netics. Further, due to the continuous segregation of iron through the sur-
face of the specimen and due to the decrease in oxygen content at the
coating's surface as a result of the volume growth of the Fe2O3 layer, a References
continuous and consistent Fe3O4 layer starts forming in this zone, with [1] M. Cingi, F. Arisoy, G. Basman, K. Sesen, Mater. Lett. 55 (2002) 360.
small uniformly distributed amounts of the dark phase. [2] L. Luo, S. Liu, J. Li, W. Yucheng, Surf. Coat. Technol. 205 (2011) 3411.
[3] R. Vaßen, M.O. Jarligo, T. Steinke, D.E. Mack, D. Stöver, Surf. Coat. Technol. 205
(2010) 938.
4. Conclusion
[4] C. Guo, J. Zhou, J. Chen, J. Zhao, Y. Yu, H. Zhou, Wear 270 (2011) 492.
[5] K. Gurumoorthy, M. Kamaraj, K.P. Rao, A.S. Rao, S. Venugopal, Mater. Sci. Eng., A
This investigation concerned the effect of PTA current variation 456 (2007) 11.
on the oxidation behavior in air at 800 and 900 °C of coatings of a [6] W. Li, Y. Li, C. Sun, Z. Hu, T. Liang, W. Lai, J. Alloys Compd. 506 (2010) 77.
[7] A. Gatto, E. Bassoli, M. Fornari, Surf. Coat. Technol. 187 (2004) 265.
nickel-based hardfacing alloy deposited on gray cast iron using two dif- [8] K. Siva, N. Murugan, R. Logesh, Int. J. Adv. Manuf. Technol. 41 (2009) 24.
ferent arc currents (100 and 128 A). Samples were analyzed by thermal [9] F. Fernandes, B. Lopes, A. Cavaleiro, A. Ramalho, A. Loureiro, Surf. Coat. Technol.
gravimetric analysis, x-ray diffraction (XRD) and scanning electron mi- 205 (2011) 4094.
[10] V. Balasubramanian, A.K. Lakshminarayanan, R. Varahamoorthy, S. Babu, J. Iron
croscopy with energy dispersion spectrometry analysis (SEM-EDS). The Steel Res. Int. 16 (2009) 44.
thermo gravimetric results show that increasing the arc current from [11] C.T. Liu, J. Ma, X.F. Sun, J. Alloys Compd. 491 (2010) 522.
100 to 128 A decreases the oxidation resistance of coatings, a result [12] W.Z. Li, Y.Q. Li, Q.M. Wang, C. Sun, X. Jiang, Corros. Sci. 52 (2010) 1753.
[13] K.-c. Zhou, L. Ma, Z.-y. Li, Trans. Nonferrous Met. Soc. China 21 (2011) 1052.
that can be correlated with the different chemical compositions of the [14] G.R. Wallwlork, Rep. Prog. Phys. 39 (1976) 401.
coatings, promoted by different dilutions of the base material. The [15] E.O. Ezugwu, Z.M. Wang, A.R. Machado, J. Mater. Process. Technol. 86 (1998) 1.
coating with lower dilution follows a parabolic oxidation weight gain [16] X. Guo, K. Kusabiraki, S. Saji, Oxid. Met. 58 (2002) 589.
[17] A.R. Lashin, O. Schneeweiss, M. Svoboda, Oxid. Met. 69 (2008) 359.
as a function of time, as this behaviour is essentially controlled by the
[18] S. Musić, S. Popović, S. Dalipi, J. Mater. Sci. 28 (1993) 1793.
growth of a Si–O layer. The coating with higher dilution follows the [19] Y.Q. Li, T. Qiu, Mater. Sci. Eng., A 444 (2007) 184.
same trend at 800 °C. However, at 900 °C two stages in the oxidation [20] V.A. Lavrenko, Y.G. Gogotsi, Oxid. Met. 29 (1988) 193.
[21] Y. Wu, Y. Niu, Scr. Mater. 53 (2005) 1247.
curve were detected, the first with a linear increase in mass gain, a com-
[22] H.K. Lee, S. Zerbetto, P. Colombo, C.G. Pantano, Ceram. Int. 36 (2010) 1589.
promise between outward Fe diffusion through a growing layer of [23] I.E. Gönenli, T.A. C., Powder Diffract. 15 (2000) 104.
Fe2O3 and the loss of carbon by decarburization and formation of CO2. [24] V.B. Trindade, B. Rodrigo, B.Z. Hanjari, S. Yang, U. Krupp, C. Hans-Jürgen, Mater.
The second step obeys a parabolic law from the moment that C libera- Res. 8 (2005) 365.
[25] Y. Nakamura, Metall. Mater. Trans. A 6 (1975) 2217.
tion is impeded by a thickened scale consisting of an external Fe2O3 [26] D.L. Douglass, J.S. Armijo, Oxid. Met. 2 (1970) 207.
phase, with small amounts of Fe3O4 and NiFe2O4 spinel rich film, and
Annex C
Annex C
F. Fernandes, A. Ramalho, A. Loureiro, A. Cavaleiro, Wear resistance of a nickel-
based coating deposited by PTA on grey cast iron, International Journal of Surface
Science and Engineering, 6 (2012) 201-213.
97
Filipe Daniel Fernandes
Int. J. Surface Science and Engineering, Vol. 6, No. 3, 2012 201
Filipe Fernandes*
CEMUC – Department of Mechanical Engineering,
University of Coimbra,
Rua Luís Reis Santos, 3030-788 Coimbra, Portugal
and
Intermolde, Rua de Leiria, 95, Apartado 103,
2431-902 Marinha Grande, Portugal
Fax: +(351) 239-790-701
E-mail: filipe.fernandes@dem.uc.pt
*Corresponding author
Abstract: The moulds for the production of glass bottles, made of cast iron, are
subjected to very severe conditions of wear during use. Thus, it is essential to
understand the wear mechanisms involved, in order to increase the equipment
life. The present study intends to study the abrasion resistance of moulds
submitted for long time at working conditions. The investigation was
conducted on nickel-based coatings deposited by plasma transferred arc (PTA)
on grey cast iron, using different arc currents. Micro-scale ball cratering
abrasive wear test was used to evaluate the tribological properties of the
as-deposited and heat treated coatings. The results show that micro-scale
abrasion tests induce grooving and pitting wear mechanisms. Increasing the arc
current decreased the hardness of the coatings and, consequently, their wear
resistance. The hardness and wear resistance of the coatings was improved by
increasing heat treatment holding time.
Keywords: plasma transferred arc; PTA; Ni-based alloy; wear; ageing; glass
moulding; surface engineering.
1 Introduction
Cast iron is commonly used in the production of moulds and accessories for glass
industry. Wear and oxidation are the main failure mechanisms that limit the surface life
of moulds (Cingi et al., 2002). Therefore, it is necessary to protect mould surfaces with
hard and wear resistant materials.
The plasma transferred arc (PTA) process has been usually applied to coat the
surfaces of moulds and, especially, their edges where the wear is higher. It produces very
high quality deposits, offering optimal protection with minimal thermal distortion of the
parts, and provides high deposition rate (Gatto et al., 2004). It allows very precise deposit
layers of complex alloys on mechanical parts that are subjected to harsh environments,
significantly extending their service life (Siva et al., 2009). However, the hardness and
the wear resistance of the coatings are strongly influenced by the dilution of the substrate
promoted by the PTA process. Low dilution provides coatings with similar chemical
composition to the metal powder added, condition to have an improved wear and
Wear resistance of a nickel-based coating deposited by PTA on grey cast iron 203
2 Experimental procedure
A nickel-based hard-facing alloy named Colmonoy 215 (from Colmonoy Company) was
deposited on flat surfaces of grey cast iron blocks by PTA process. Before coating, each
cast iron block was preheated at 480°C, in order to reduce the thermal shock and cooling
rate to avoid cracks in both the coating and heat-affected zone. The same process
204 F. Fernandes et al.
parameters were used for all coatings except arc current, which was 100 A for
specimen 1, 128 A for specimen 2 and 140 A for specimen 3. The composition of cast
iron and metal powder is given in Table 1 and the process parameters used in Table 2.
After coating, the blocks were left to cool to room temperature. Samples were removed
from each block, transversely to welding direction, for further analysis. Ageing studies
were carried out at 500°C for 5, 10 and 20 days on specimen 3.
Table 1 Nominal chemical composition (wt.%) of substrate and hard-facing alloy
Base material C Mn Si P S Cr Ni Mo V Ti Fe
Grey cast iron 3.60 0.60 2.00 < 0.20 < 0.04 < 0.20 < 0.50 0.50 0.10 0.20 Balance
Hard-facing C Cr Si B Fe Al F Co Ni
Ni-alloy 0.14 2.45 2.56 0.86 1.08 1.30 0.01 0.08 Balance
Main arc current Powder feed rate Travel speed Powder feed gas
(A) (rpm) (mm/s) flow rate (l/min)
Specimen 1 100 20 2 2
Specimen 2 128 20 2 2
Specimen 3 140 20 2 2
Plasma gas Shielding gas Torch work Preheat
Oscillation
flow rate flow rate distance temperature
(mm)
(l/min) (l/min) (mm) °C
Specimen 1 2.2 20 13 4 480
Specimen 2 2.2 20 13 4 480
Specimen 3 2.2 20 13 4 480
Microstructural analysis was done in cross section of the samples. Optical and scanning
electron microscopies (OM and SEM) were used to characterise the microstructure of the
coatings. Vickers micro-hardness profiles were measured in the cross section of treated
samples using a load of 5 N.
Micro abrasive wear tests were carried out in as-deposited and heat treated surfaces at
room temperature in a ball cratering devices. Figure 1 shows a schematic diagram of the
equipment used. In this test a specific normal load is applied to press a rotating ball into
the flat surface of the coated sample in the presence of a slurry suspension of abrasive.
The abrasive slurry, agitated by a magnetic stirrer, was continuous and gravitationally
drip fed onto the rotating ball to produce a spherical cup depression. Preliminary tests
using different contact conditions were done to reproduce the interaction conditions in the
surface of the moulds which have been in service for a long time. The wear mechanism
was studied by SEM. High purity silica SS40 with angular particles of averaging of
3.10 μm in size was used as an abrasive to replicate the effect of the melted glass. The
slurry used was prepared with a concentration of 103 g per 1 × 10–4 m3 of distilled water.
A steel ball bearing (AISI 52100) with 0.0254 m was used. Several scars were made in
each sample with different durations, 100, 200, 300, 400 and 500 turns, respectively 7.98,
15.96, 23.94, 31.92, 39.90 meters of sliding distance for all the specimens, under the
action of a normal load of 0.1 N. In order to minimise the scatter of results, previously to
starting the tests, the ball surface was etched with chloridric acid (33%) during ten
Wear resistance of a nickel-based coating deposited by PTA on grey cast iron 205
minutes. The dimensions of spherical depressions were measured in order to calculate the
wear volume, which is given by the equation (1). In this equation, R represents the ball
radius and b the crater chordal diameter of the spherical cup depression. The b value was
evaluated using the average of two perpendicular measurements performed at the wear
scar chordal diameter. Those values were measured with a Mitutoyo Toolmaker
Microscope with x-y micrometer table.
V = ( π × b 4 ) × (64 × R) (1)
In order to obtain the specific wear rate, to quantify the wear behaviour, Archard’s law
was applied, see equation (2). This is the most applied law to determine a parameter to
quantify the wear behaviour of materials. In this equation, P is the normal load, l is the
sliding distance and k is the specific wear rate.
V = k×P×l (2)
To quantify the specific wear rate a linear approach of Archard’s model was applied and
a statistical analysis was used to estimate the error of measurements, as described
elsewhere (Ramalho, 2010).
3 Results
3.1 Microstructure
Figure 2 shows the microstructure of the coatings performed using different arc currents
[Figures 2(a), 2(b) and 2(c)], as well as the fusion boundary of specimen 3 [Figure 2(d)],
and SEM pictures of specimen 3 before and after heat treatment [Figures 2 (e) and 2(f)].
The typical microstructure of deposits consists of dendrites of Ni-Fe solid solution phase
with columnar morphology oriented along the direction of heat flow, with C-flakes (dark
floret-like structures) randomly distributed. However, near the interface it was observed
the presence of a cellular microstructure finer than the dendrite structure with diminution
of the size of C-flakes, as illustrated in Figure 2(d) for specimen 3. This behaviour could
be explained by the higher solidification rates involved in the boundary because of the
206 F. Fernandes et al.
efficient thermal exchange ensured by the high volume ratio substrate/coating of the
base material, as suggested by Gatto et al. (2004) and Navas et al. (2006). Moreover, the
presence of the ‘dark’ phase on the coatings cannot be explained by the chemical
composition of the metal powder added, so it can be only attributed to the high dilution of
cast iron induced by the PTA process. The dilution was evaluated by the ratio between
the area of the melted base material and the total area of the melted zone, both measured
in the cross section of the coatings. Dilutions of 28%, 50% and 59% were measured
respectively for specimen 1, 2 and 3. These results agree with the observed increasing
proportion of C-flakes for higher arc currents, as can be seen in Figures 2(a), 2(b) and
2(c). Therefore, the chemical composition of the coatings is largely influenced by the
dilution of cast iron. A detailed characterisation of this type of coatings can be found in a
previous publication of the authors (Fernandes et al., 2011).
Figure 2 Microstructure of (a) specimen 1, (b) specimen 2, (c) specimen 3, (d) interface of
specimen 3 and SEM pictures of sample 3 (e) in non-heat-treated condition and
(f) in heat treated condition during ten days
Wear resistance of a nickel-based coating deposited by PTA on grey cast iron 207
Table 3 SEM chemical analysis of the deposit and base material of specimen 3
Elements wt%
C Al Si Cr Fe Ni
1 0.8 0.7 2.6 1.4 54.7 39.8 Deposit
2 0.8 0.7 2.8 1.6 54.8 39.2
3 0.9 1.4 2.5 1.9 50.4 42.8
4 1.6 1.1 2.2 1.4 59.4 34.2 Interface
5 3.9 - 2.7 - 93.3 - Base material
Figure 3 SEM magnification of microstructure of the coating 3 (a) in non-heat treated condition
and (b) in heat-treated condition after ten days exposed at 500°C and SEM EDAX
spectra of (c) point 1, (d) point 2, and (e) point 3 (see online version for colours)
(c) (d)
(e)
Table 3 shows the chemical composition in area measured by SEM-EDS at various zones
on the coating and the substrate of specimen 3. Zone 1 and 2 are located close to the
coating surface, zone 3 at the mid-thickness of the coating, zone 4 in the coating close to
the fusion boundary and zone 5 in the bulk cast iron. This analysis showed that PTA
process promotes an evenly distribution of the elements inside the coating. The same
observation was done in specimens 1 and 2. Due to the dilution, the amount of iron in the
208 F. Fernandes et al.
coating increases reducing its nickel content. Ezugwu et al. (1998) reported that the
increase of iron content in the coatings tends to decrease their oxidation resistance,
because of the formation of a less adherent oxide scale.
Figures 2(e) and 2(f) reveals that the heat treatment induces significant changes in the
original microstructure of specimen 3. It promotes the formation of a white phase on the
grain boundaries. Magnifications of the two microstructures are shown in Figure 3.
SEM-EDS spectra analyses were conducted on the coatings to give a detailed
characterisation of their microstructure. From the EDS analysis, it can be stated that the
grain boundaries of the non-heat-treated sample is formed by two phases: a light grey
and a dark grey rich in chromium and silicon, respectively. Inversely, the specimen
in heat-treated condition is formed by the light grey phase as observed in the
non-heat-treated sample and a new phase (white phase) which is replacing the dark phase.
The amount of the white phase grows with increasing time of heat treatment. According
to Fernandes et al. (2011), who studied a nickel alloy deposited on grey cast iron, the
dark grey phase is a Ni3Si phase. The EDS spectrum reveals that both phases (the dark
and the white) are rich in Ni-Si. However, the white phase presents a higher content of
phosphorous on the heat treated sample suggesting a dispersion of Ni-P precipitates on
the grain boundaries, as observed by Sahoo and Das (2011) in electroless nickel coatings.
Figure 4 Hardness profile across the interface of (a) as-deposited specimens and (b) specimen 3
before and after heat treatment (see online version for colours)
(a) (b)
3.2 Hardness
The micro-hardness profiles of the different as-deposited and heat treated specimens are
shown respectively in Figures 4(a) and 4(b). Figure 4(a) reveals that the hardness through
the coatings is approximately constant and tends to decrease with increasing arc current.
The dilution induced by the PTA process can explain this behaviour. The partial melted
zone (PMZ) displays the highest hardness. According to Pouranvari (2010), who studied
the weldability of grey cast iron using nickel-based filler metals, the highest hardness
observed in PMZ is due to the formation of hard and brittle phases during surfacing, such
as martensite and white cast iron. The heat treatment promotes the progressive hardening
of the coatings with increasing thermal exposure time [see Figure 4(b)]. Some scatter of
the hardness values can be observed too. Harsha et al. (2008), who studied the influence
of CrC in a nickel flame sprayed coatings, observed similar behaviour. According to
them, the scatter in the values is due to the presence of a harder eutectic network around
the cells. As described before, heat treatment promotes the formation of Ni-P precipitates
Wear resistance of a nickel-based coating deposited by PTA on grey cast iron 209
on the grain boundaries. The greater hardness observed on the ageing treated samples can
be attributed to the presence of these precipitates similarly to what is currently observed
in Ni-P electroless coating when annealed at increasing temperatures (Sahoo and Das,
2011). This type of coatings is a supersaturated alloy in the as-deposited state and can be
strengthened by the precipitation of nickel phosphide crystallites. The phosphides act as
barriers for dislocation movement and improve the mechanical strength of the coatings.
Figure 5 Evolution of the wear volume as function of the parameter ‘normal load × sliding
distance’ of (a) as-deposited coatings and (b) specimen 3 before and after thermo
exposure at 500°C during several days (see online version for colours)
(a) (b)
The results of the micro-scale abrasion tests from specimens in as-deposited and heat
treaded conditions are plotted respectively in Figures 5(a) and 5(b), as the material
volume loss as a function of the product of the sliding distance (l) by the normal load (P).
The reliability analysis of the data is presented in Table 4 and the confidence interval for
the specific wear rate plotted in Figure 6. For each sample, a straight linear trend was
fitted to the experimental points in order to obtain the specific wear rate, which
corresponds to its slope [see equation (2)]. From Figure 5(a), it is possible to conclude
that specimen 1 exhibits slightly lower specific wear rate, therefore it is the more resistant
to wear. This result is in accordance with its greater hardness when compared to the
others as-deposited coatings. Specimens 2 and 3 displays similar wear volumes, which is
210 F. Fernandes et al.
again attributed to their similar hardness profiles. Figure 5(b) reveals that specimens
exposed during five and ten days at 500°C displayed similar specific wear rate to the
non-heat-treated specimen. However, the specimen exposed during 20 days exhibits
lower specific wear rate, therefore, it is the most resistant to wear. This last result is again
in accordance with the higher hardness of this sample after the heat treatment.
Table 4 Results of the linearisation analysis of samples in as-deposited and heat treated
conditions
Figure 6 Confidence interval for a confidence level of 90%, for the specific wear rate of
specimens in as-deposited and heat treated conditions (see online version for colours)
In order to identify the wear mechanisms and to compare the worn surfaces with those
observed in the surface of moulds, SEM examinations were carried out. Figure 7(a)
displays a SEM micrograph of the surface of a coating of a mould after long time in
service. From this picture is possible to see that the surface failure regions display pits
and cracks. Considering the in-service conditions the surface damage probably was
induced by synergetic effects of several mechanisms, namely thermal fatigue, corrosion
and abrasion. However, pitting is the main active failure mechanism of the moulds.
Figures 7(b) and 7(c) illustrates spherical cup depressions induced by the micro-scale
abrasion tests on the specimen 3, before and after holding time for 20 days, produced
after 500 ball rotations. Both scars display similar wear behaviour with two modes of
wear: three-body abrasion (pitting) and two-body abrasion (grooving). The two-body
abrasion mode is preferentially localised in the area of higher pressure (input slurry area)
Wear resistance of a nickel-based coating deposited by PTA on grey cast iron 211
and three-body abrasion mode is localised in the area of less pressure. Figure 8 shows
these two modes of wear at higher magnification. The same wear mechanisms were
observed in the other specimens in as-deposit and heat treated conditions. Although the
scars displayed two modes of wear it can be stated that the region corresponding to
three-body mode reproduces closely the wear mechanism present in the moulds surface.
However, one should keep in mind that these results do not reproduce with accuracy the
in-service condition of the moulds, since in service they are exposed to other damage
mechanisms such as thermal and mechanical fatigue, oxidation and corrosion at high
temperature.
Figure 7 SEM micrograph (a) of the surface of a coating of a mould after long exposure to high
temperature, (b) scar with 500 rotations induced by the micro-scale abrasion tests on
specimen 3, and (c) scar with 500 rotations induced by the micro-scale abrasion tests on
specimen 3 after thermal exposure for 20 days (see online version for colours)
Figure 8 SEM micrograph of (a) three-body abrasion and (b) two-body abrasion
212 F. Fernandes et al.
4 Conclusions
Micro-scale abrasion testing was used to study the interaction conditions between melted
glass and the surface of the moulds. The influence of the arc current used in PTA process,
as well as the effect of heat treatments, on the microstructure, hardness and wear
resistance of the coatings was analysed. The increase in arc current increased the dilution
of the base material changing the composition and microstructure of the deposits and
reducing their hardness. Furthermore, the wear loss of material increased with increasing
arc current. The micro-scale abrasion tests displayed two modes of wear: three-body
abrasion (pitting) and two-body abrasion (grooving). Although the scars displayed two
modes of wear it can be stated that the wear mechanism of the abrasion test reproduces
with success the wear mechanism present in the moulds surface. The heat treatment
performed on the sample 3 promotes increasing hardness of the coating with thermal
exposure time. The enhanced hardness of the coatings is partially attributed to Ni-P
precipitates on the grain boundaries. Specimen 3 heat treated for 20 days showed the best
resistance to wear.
Future experiments will be devoted to study the high temperature abrasion resistance
of the coatings, in order to reproduce the in service conditions of moulds.
Acknowledgements
The authors would like to thank the company ‘Intermolde’, for supplying the coated
samples, and the Portuguese Foundation for the Science and Technology (FCT), through
COMPETE programme from QREN and to FEDER, for financial support in the aim of
the project number ‘013545’, as well as for the grant SFRH/BD/68740/2010.
References
Balasubramanian, V., Lakshminarayanan, A.K., Varahamoorthy, R. and Babu, S. (2009)
‘Application of response surface methodolody to prediction of dilution in plasma transferred
arc hardfacing of stainless steel on carbon steel’, International Journal of Iron and Steel
Research, Vol. 16, No. 1, pp.44–53.
Chang, J.H., Chang, C.P., Chou, J.M., Hsieh, R.I. and Lee, J.L. (2010) ‘Microstructure and bonding
behavior on the interface of an induction-melted Ni-based alloy coating and AISI 4140 steel
substrate’, Surface and Coatings Technology, Vol. 204, No. 2, pp.3173–3181.
Cingi, M., Arisoy, F., Basman, G. and Sesen, K. (2002) ‘The effects of metallurgical structures of
different alloyed glass mold cast irons on the mold performance’, Materials Letters, Vol. 55,
No. 6, pp.360–363.
Ezugwu, E.O., Wang, Z.M. and Machado, A.R. (1998) ‘The machinability of nickel-based alloys: a
review’, Journal of Materials Processing Technology, Vol. 86, Nos. 1–3, pp.1–16.
Fernandes, F., Lopes, B., Cavaleiro, A., Ramalho, A. and Loureiro, A. (2011) ‘Effect of arc current
on microstructure and wear characteristics of a Ni-based coating deposited by PTA on gray
cast iron’, Surface and Coatings Technology, Vol. 205, No. 16, pp.4094–4106.
Gant, A.J. and Gee, M.G. (2011) ‘A review of micro-scale abrasion testing’, Journal of Physics D:
Applied Physics, Vol. 44, No. 7, 073001.
Gatto, A., Bassoli, E. and Fornari, M. (2004) ‘Plasma transferred arc deposition of powdered high
performances alloys: process parameters optimisation as a function of alloy and geometrical
configuration’, Surface and Coatings Technology, Vol. 187, Nos. 2–3, pp.265–271.
Wear resistance of a nickel-based coating deposited by PTA on grey cast iron 213
Guo, C., Zhou, J., Chen, J., Zhao, J., Yu, Y. and Zhou, H. (2011) ‘High temperature wear resistance
of laser cladding NiCrBSi and NiCrBSi/WC-Ni composite coatings’, Wear, Vol. 270,
Nos. 7–8, pp.492–498.
Gurumoorthy, K., Kamaraj, M., Rao, K.P., Rao, A.S. and Venugopal, S. (2007) ‘Microstructural
aspects of plasma transferred arc surfaced Ni-based hardfacing alloy’, Materials Science and
Engineering: A, Vol. 456, Nos. 1–2, pp.11–19.
Harsha, S., Dwivedi, D. and Agarwal, A. (2008) ‘Influence of CrC addition in Ni-Cr-Si-B flame
sprayed coatings on microstructure, microhardness and wear behaviour’, The International
Journal of Advanced Manufacturing Technology, Vol. 38, No. 1, pp.93–101.
Kesavan, D. and Kamaraj, M. (2010) ‘The microstructure and high temperature wear
performance of a nickel base hardfaced coating’, Surface and Coatings Technology, Vol. 204,
No. 24, pp.4034–4043.
Navas, C., Colaço, R., Damborenea, J. and Vilar, R. (2006) ‘Abrasive wear behaviour of laser clad
and flame sprayed-melted NiCrBSi coatings’, Surface and Coatings Technology, Vol. 200,
No. 24, pp.6854–6862.
Pouranvari, M. (2010) ‘On the weldability of grey cast iron using nickel based filler metal’,
Materials & Design, Vol. 31, No. 7, pp.3253–3258.
Ramalho, A. (2010) ‘A reliability model for friction and wear experimental data’, Wear, Vol. 269,
Nos. 3–4, pp.213–223.
Ramalho, A., Rodríguez, J. and Rico, A. (2011) ‘Microabrasion resistance of nanostructured
plasma-sprayed coatings’, Int. J. Surface Science and Engineering, Vol. 5, No. 4, pp.250–260.
Sahoo, P. and Das, S.K. (2011) ‘Tribology of electroless nickel coatings – a review’, Materials &
Design, Vol. 32, No. 4, pp.1760–1775.
Siva, K., Murugan, N. and Logesh, R. (2009) ‘Optimization of weld bead geometry in plasma
transferred arc hardfaced austenitic stainless steel plates using genetic algorithm’, The
International Journal of Advanced Manufacturing Technology, Vol. 41, No. 1, pp.24–30.
Annex D
Annex D
F. Fernandes, A. Ramalho, A. Loureiro, A. Cavaleiro, Mapping the micro-
abrasion resistance of a Ni-based coating deposited by PTA on gray cast iron, Wear,
292-293 (2012) 151-158.
113
Filipe Daniel Fernandes
Wear 292–293 (2012) 151–158
Wear
journal homepage: www.elsevier.com/locate/wear
a r t i c l e i n f o abstract
Article history: Micro-scale abrasion was used to characterize abrasion resistance of a nickel-based hardfacing alloy
Received 2 February 2012 deposited by Plasma Transferred Arc on gray cast iron using silica as abrasive agent. In order to
Received in revised form investigate the occurrence of different abrasion mechanisms, several test conditions were used,
28 May 2012
namely: different silica abrasive contents and a range of normal loads and test durations. A scanning
Accepted 30 May 2012
Available online 7 June 2012
electron microscope (SEM) was used to study the morphologies of the spherical cup-shaped depres-
sions induced by different test conditions. The results are discussed in terms of the effect of the
Keywords: dominant wear mechanisms on the abrasion resistance and the influence of the test conditions on the
Mapping mechanism transition. The wear results lead to the conclusion that the specific wear rate essentially
Micro-scale abrasion
depends on the wear mechanisms (rolling or grooving) involved and not on the tests conditions
Three-body abrasion
employed, since these do not produce changes in the wear mechanism.
Two-body abrasion
Hardfacing & 2012 Elsevier B.V. All rights reserved.
Wear testing
0043-1648/$ - see front matter & 2012 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.wear.2012.05.018
152 F. Fernandes et al. / Wear 292–293 (2012) 151–158
3. Results
Fig. 3. Silica abrasive particles. (a) Morphology and (b) particle size distribution.
Table 1
Test conditions.
spherical cup depression produced and R is the ball radius used. The
b value was calculated using the average of two perpendicular
measurements performed at the wear scar chordal diameter, with
the help of a Mitutoyo Toolmaker Microscope with x–y micrometer
table. All the data was then collected in a single graph, and
represented as function of the volume loss, with the product
between the sliding distance and the normal load, to identify the
conditions that produce similar wear mechanisms. To quantify the
specific wear rate of the material a linear approach to Archard’s law
was used. This law is given by Eq. (2). In this equation V represents
the wear volume, k is the specific wear rate, N is the normal applied
load and x is the sliding distance. Moreover, a statistical analysis was
applied to estimate the error of measurements, as described else-
where [22].
4
V ¼ ðf b Þ=ð64 RÞ ð1Þ
V ¼kNx ð2Þ
stabilizes with increasing sliding distance, whatever the concen- analyses were conducted on different spherical cup depressions in
tration of abrasive used. This behavior could be associated with order to identify the wear mechanism produced by the dissimilar
two effects: the running-in effect [23] and the variation of the test conditions.
contact area/wear volume ratio throughout the duration of the Fig. 5 shows SEM micrographs of the dissimilar morphologies
test. According to Blau [23], the running-in behavior is the net of the spherical cup depressions induced by different test condi-
result of simultaneous transitional processes occurring within the tions of the micro-scale abrasion tests. As these images show, two
interface, which may lead to non-linear evolution of the wear for different wear mechanisms were identified, rolling abrasion (see
short sliding distances. Furthermore, at the start of the wear tests Fig. 5(a)) and grooving abrasion (see Fig. 5(b)). Moreover, a
(with very small craters, consequently with small amounts of the combination of both wear mechanisms (see Fig. 5(c)) was also
radius of the spherical depression) it is well known that the area identified for several contact conditions. In order to facilitate
of the crater grows much faster than the volume; at that point, identification of the wear mechanisms present in each worn
the rate of pressure reduction is greater than the rate of change in surface, magnified images are shown in Fig. 5(a1)–(c1). Rolling
the volume, which may also justify the reduction in wear rate. In abrasion is characterized by multiple indentations in the worn
fact, the severity of contact index, proposed by Adachi and surface, (see Fig. 5(a1), whereas grooving abrasion (see Fig. 5(b1)
Hutchings [6], is inversely proportional to the crater area; there- is characterized by parallel grooves in the worn surface which
fore, a sudden increase in the area results in a reduction of the results from the plastic deformation of the material. In turn, the
severity index, which will induce a change in the wear rate. The surface morphology, which displays a mixed-mode (rolling and
changes in the wear volume described above are only partially grooving abrasion), corresponds to a transition zone characterized
correlated with the test conditions, encouraging the investigation by a mixture of grooves and indentations, as shown in Fig. 5(c1).
of the interactions between the abrasive slurry and the surface of The transition of rolling to grooving abrasion occurs due to the
the specimen in the contact zone. Scanning electron microscopy embedding of particles in the surface. In this case grooving starts
Fig. 5. SEM morphologies of the worm surfaces of the sample performed with the combination of load, volume fraction of abrasive and sliding distance of: (a) 0.5 N, 50%
and 4 m, (b) 0.5 N, 50%, 16 m and (c) 0.1 N, 25%, 16 m. Magnification of the wear mechanism present in the wear scar of: (a1) Figure a), (b1) Figure b), (c1) Figure c).
F. Fernandes et al. / Wear 292–293 (2012) 151–158 155
Fig. 6. (a) Scanning electron micrograph of the wear scar produced with 25% of
volume fraction of abrasive, load of 0.5 N and a sliding distance of 6.4 m: 1–ridging
wear mechanism zone, 2–grooving wear mechanism zone and 3–rolling wear
mechanism zone. (b) Magnification of the ridging wear mechanism zone.
these results, the specific wear rate can be estimated using points
produced with different test conditions, since a single wear mode
is sustained. Furthermore, these results lead to the conclusion
that the specific wear rate essentially depends on the wear
mechanisms (rolling or grooving) involved and not on the tests
conditions employed, since these do not produce changes in the
wear mechanism.
5. Conclusion
(
1 ( ðb=RÞ r0:345
CF ¼ ^
1=ð10:66 ð1ð1ðb=ð2 RÞÞ2 Þð1=2Þ ÞÞ ( 0:345 rðb=RÞ r1:786
ðA6Þ
References
[1] K. Bose, R.J.K. Wood, Optimum tests conditions for attaining uniform rolling
abrasion in ball cratering tests on hard coatings, Wear 258 (2005) 322–332.
[2] A.J. Gant, M.G. Gee, A review of micro-scale abrasion testing, Journal of
Physics D: Applied Physics 44 (2011) 073001.
[3] M.F.C. Andrade, R.P. Martinho, F.J.G. Silva, R.J.D. Alexandre, A.P.M. Baptista,
Influence of the abrasive particles size in the micro-abrasion wear tests of
Fig. A.2. Error associated to the approximate equation to estimate the volume of a TiAlSiN thin coatings, Wear 267 (2009) 12–18.
wear scar without and with corrective factor.. [4] T.S. Eyre, Wear characteristics of metals, Tribology International 9 (1976)
203–212.
[5] D. Braga, A. Ramalho, P.N. Silva, A. Cavaleiro, Study of abrasion resistance of
equation can be calculated by using Eq. (A3). steels by micro-scale tests, in: P.M. Vilarinho (Ed.), Advanced Materials
Forum III, Parts 1 and 2, Trans Tech Publications Ltd., Zurich-Uetikon, 2006,
2 pp. 544–548.
V ¼ ð1=6Þ f h ð3 ðb=2Þ2 þ h Þ ðA1Þ
[6] K. Adachi, I.M. Hutchings, Wear-mode mapping for the micro-scale abrasion
test, Wear 255 (2003) 23–29.
4
V n ¼ ðf b Þ=ð64 RÞ ðA2Þ [7] R.C. Cozza, D.K. Tanaka, R.M. Souza, Friction coefficient and wear mode
transition in micro-scale abrasion tests, Tribology International 44 (2011)
1878–1889.
error ¼ ðV n VÞ=V ðA3Þ [8] M.M. Stack, M. Mathew, Micro-abrasion transitions of metallic materials,
Wear 255 (2003) 14–22.
It is evident from Fig A.1 that increasing the depth h of the
[9] R.J.K. Wood, Tribo-corrosion of coatings: a review, Journal of Physics D:
spherical cup depression increases the value of the chordal Applied Physics 40 (2007) 5502–5521.
diameter b of the wear scar, until a maximum value of b equal [10] R.C. Cozza, J.D.B. de Mello, D.K. Tanaka, R.M. Souza, Relationship between test
severity and wear mode transition in micro-abrasive wear tests, Wear 263
to 2R is reached. So, discretizing the value of b in small intervals it
(2007) 111–116.
is possible to calculate the associated values of V and Vn for each [11] E. Rabinowicz, L.A. Dunn, P.G. Russell, A study of abrasive wear under three-
relation between b and h. The value of h (given by Eq. (A4)), can be body conditions, Wear 4 (1961) 345–355.
easily deduced by subtracting the value of c from the value of R, as [12] J.C.A. Batista, A. Matthews, C. Godoy, Micro-abrasive wear of PVD duplex and
single-layered coatings, Surface and Coatings Technology 142–144 (2001)
represented in Fig A.1. The value of c is obtained by applying 1137–1143.
Pythagoras’ Theorem at the triangle. Applying Eq. (A3) at the [13] J.A. Williams, A.M. Hyncica, Mechanisms of abrasive wear in lubricated
dissimilar combinations of b and h it is possible obtain the error contacts, Wear 152 (1992) 57–74.
[14] J.A. Williams, A.M. Hyncica, Abrasive wear in lubricated contacts, Journal of
associated by using an approximate equation to estimate the Physics D: Applied Physics 25 (1992) A81–A90.
volume of a spherical cup depression induced by the micro-scale [15] F.J.G. Silva, R.B. Casais, R.P. Martinho, A.P.M. Baptista, Role of abrasive
abrasion test. The results of the error are plotted in Fig A.2, dashed material on micro-abrasion wear tests, Wear 271 (2011) 2632–2639.
[16] R.I. Trezona, D.N. Allsopp, I.M. Hutchings, Transitions between two-body and
line, as function of the ratio between the crater’s chordal diameter three-body abrasive wear: influence of test conditions in the microscale
and the radius of the ball with the relative error associated. As can abrasive wear test, Wear 225–229 (1) (1999) 205–214.
be seen in this figure, the resulting error is very small for low [17] M.M. Stack, M.T. Mathew, Mapping the micro-abrasion resistance of WC/Co
based coatings in aqueous conditions, Surface and Coatings Technology 183
ratios of b/R, however with an increase in this ratio, the error
(2004) 337–346.
increases. For example considering a spherical cup depression [18] A. Ramalho, J.R. Pérez, Á.R. Garcia, Microabrasion resistance of nanostruc-
with a chordal diameter equal to the radius of the ball, the error tured plasma-sprayed coatings, International Journal of Surface Science and
Engineering 5 (2011) 250–260.
committed in the calculations is about 10 percent.
[19] P.H. Shipway, J.J. Hogg, Wear of bulk ceramics in micro-scale abrasion—the
role of abrasive shape and hardness and its relevance to testing of ceramic
h ¼ RðR2 ðb=2Þ2 Þ1=2 ðA4Þ
coatings, Wear 263 (2007) 887–895.
The level of correlation between the ratios of the depth of the [20] F. Fernandes, B. Lopes, A. Cavaleiro, A. Ramalho, A. Loureiro, Effect of arc
current on microstructure and wear characteristics of a Ni-based coating
spherical cup depression and the ball radius with the error proved to deposited by PTA on gray cast iron, Surface and Coatings Technology 205
be linear. So, according to this relationship it is possible to establish a (2011) 4094–4106.
158 F. Fernandes et al. / Wear 292–293 (2012) 151–158
[21] M.G. Gee, A.J. Gant, I.M. Hutchings, Y. Kusano, K. Schiffman, K.V. Acker, [23] P.J. Blau, On the nature of running-in, Tribology International 38 (2005)
S. Poulat, Y. Gachon, J.v. Stebut, P. Hatto, G. Plint, Results from an inter- 1007–1012.
laboratory exercise to validate the micro-scale abrasion test, Wear 259 [24] K. Adachi, I.M. Hutchings, Sensitivity of wear rates in the micro-scale
(2005) 27–35. abrasion test to test conditions and material hardness, Wear 258 (2005)
[22] A. Ramalho, A reliability model for friction and wear experimental data, Wear 318–321.
269 (2010) 213–223.
Annex E
Annex E
F. Fernandes, T. Polcar, A. Loureiro, A. Cavaleiro, Room and high temperature
tribological behavior of Ni-based coatings deposited by PTA on gray cast iron,
(2014), under review, “Tribology International”.
123
Filipe Daniel Fernandes
Annex E
1
CEMUC - Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis
Santos, 3030-788 Coimbra, Portugal.
2
Department of Control Engineering Czech Technical University in Prague Technicka 2,
Prague 6, 166 27 Czech Republic.
3
n-CATS University of Southampton Highfield Campus SO17 1BJ Southampton, UK.
*
Email address: filipe.fernandes@dem.uc.pt, tel. + (351) 239 790 745, fax. + (351) 239 790
701
Abstract
In the present investigation the effect of the substrate dilution on room and high
temperature (550 and 700 ºC) tribological behavior of nickel based hardfaced coating
deposited by plasma transferred arc onto a gray cast iron was investigated and compared to
the uncoated gray cast iron. At room temperature, the wear loss of coatings was independent
of the substrate dilution and similar to the gray cast iron. At high temperatures, coating
produced with high dilution displayed the highest wear resistance between all the samples.
This is attributed to the formation of a protective tribo-layer resulting from the agglomeration
of a high amount of oxide debris due to its lower oxidation resistance when compared to the
sample produced with low dilution.
Keywords: Plasma transferred arc, Ni-based alloy, Substrate dilution effect, Tribology, High
temperature wear
1. Introduction
Cast irons and copper-alloys are commonly used in the production of glass molds and
accessories for glass industry, owing to their excellent thermal conductivity and relatively low
cost [1]. Molds are often exposed to severe conditions of abrasion, oxidation, wear, fatigue at
high temperature, due to repeated contact with melted glass, causing deformation or failure of
125
Filipe Daniel Fernandes
Annex E
parts and, thus, compromising the product quality and increasing the maintenance costs. To
overcome this shortcoming protective coatings are normally applied at molds surfaces with
the aim of increasing their lifetime in harsh environments [2, 3]. A wide diversity of coating
materials have been successfully applied to protect the surface of these components, such as:
cobalt, iron and nickel alloys [2, 4, 5]. The latter have been especially used due to their
outstanding performance under extreme high temperature conditions [5-8]. Several hardfacing
and thermal spraying processes have been used to coat the molds: (i) plasma transferred arc
(PTA) and gas tungsten arc welding (GTAW), (ii) flame spraying (FS), high velocity oxy fuel
(HVOF) and atmospheric plasma spraying (APS). Hardfacing processes have been
preferentially used in the protection of mold surfaces because they promote a strong
metallurgical bonding between the coating and the substrate, condition required to achieve
high quality adherent coatings and to avoid the catastrophic failure in service, as opposed to
thermal spraying processes where only a mechanical bonding is established. Among the
hardfacing processes referred to above, PTA has been widely used to protect the surface of
molds, since it produces high quality thick coatings, offering both optimal protection with
minimal thermal distortion of parts and high deposition rates in single layer deposits [4, 5].
However, the properties and the quality of deposits are strongly dependent on the dilution of
the substrate promoted by the PTA process. Low dilution provides coatings with similar
chemical composition to the added metal powder, condition for achieving enhanced
mechanical properties, wear and oxidation resistance [9]. To avoid adhesion problems and,
therefore, not to compromise the performance of molds in service, it is common to increase
the dilution through the change of the most relevant deposition parameters even if the
mechanical properties, wear and oxidation resistance are diminished. In our previous studies
[10, 12] the effect of increasing dilution, promoted by change of arc current on the structure,
mechanical properties, oxidation resistance and wear behavior of a nickel-based alloy
deposited onto a gray cast iron have been investigated. Regarding the wear behavior of
coatings, ball cratering wear tester has been selected to reproduce the wear mechanism
produced by the interaction between the melted glass and the surface of molds (three body
abrasion). However, this equipment does not allow consider the high temperature effect.
Hence, as a way to complement the previous studies, the aim of this investigation is to study
the effect of increasing substrate dilution on the tribological performance at room and
essentially at high temperature of as deposited Ni-based coating deposited onto gray cast iron,
using a pin on disc tribometer apparatus. Comparison of these results with those achieved for
the base material is also provided.
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2. Experimental procedure
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electron microscopy with energy dispersive X-ray spectroscopy (SEM-EDS) and by Raman
spectroscopy.
Table 2 - Nominal chemical composition (wt.%) of the gray cast iron and nickel-based alloy powder.
Base material C Mn Si P S Cr Ni Mo V Ti Fe
Gray cast iron 3.60 0.60 2.00 < 0.20 < 0.04 < 0.20 <0.50 0.50 0.10 0.20 Balance
Hardfacing C Cr Si B Fe Al F Co Ni
Ni-alloy 0.14 2.45 2.56 0.86 1.08 1.30 0.01 0.08 Balance
The typical microstructure of gray cast iron and coatings produced with 100 and 140 A
is depicted in Fig. 1. Both coatings display relatively dense microstructure with excellent
adhesion to the substrate, free of microcracks and only a few solidification voids. Their
typical microstructure consists of dendrites of Ni-Fe solid solution phase aligned along the
direction of heat flow. Flakes of carbon (dark-floret like phase) randomly distributed in the
matrix can be perceived too, being more notorious on C140 coating. The high amount of dark-
floret phase on C140 coating is in good agreement with the observed increase of dilution with
increasing arc current. Following the procedure described in reference [11], dilutions of ~28
and ~59% were measured for C100 and C140 coatings, respectively. The dilution is beneficial
in terms of improving the adhesion, but it can bring some disadvantages, such as change in
the hardness and chemical composition, and lowering of wear and oxidation resistance of the
coatings [9]. The chemical composition of coatings, assessed by SEM-EDS given in Table 3,
were evaluated from the average of several measurements performed on an area 400×400 µm
in the cross section at middle thickness of each coating. This procedure was used to ensure
that the chemical composition was determined in a representative volume of material in that
zone, thus avoiding problems that punctual analyses can give in heterogeneous materials. The
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results show that the main differences were in the iron, silicon and carbon contents, in good
agreement with the different dilutions in the samples. The gray cast iron shows flakes of
graphite evenly distributed in a ferritic matrix.
Graphite
Graphite
a) b)
Graphite
Coating
c) d) Grains of ferrite
Figure 1 - Optical micrograph of: a) C100 coating, b) C140 coating, c) interface between coating and substrate of
C140 coating, d) gray cast iron.
wt. % of elements
Coating Ni Fe C Si Al Cr
C100 85.40 8.47 0.30 2.17 0.81 2.86
C140 55.71 37.73 1.26 2.52 0.62 2.16
Fig. 2 shows the Vickers hardness profiles across the interface of the deposited
coatings, measured with 5 N applied load. As it would be expected, the increase in the weld
current, and therefore increase in dilution, gives rise to a decrease of the hardness either in the
coating, undesirable in terms of wear, or in the heat affected zone (HAZ), which is favourable
in terms of toughness of this region. The average surface hardness of coatings C100, C140
and gray cast iron was 300, 195 and 150 Vickers (HV0.5), respectively. The high hardness of
the HAZ zone in relation to the non-thermal affected based material is related to the formation
of hard and brittle phases. The detailed characterization of the microstructure of each coating
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and the hardness variations with increasing dilution can be found in a previous publication of
the authors [11].
500
Fusion Line
PMZ
450 C100 coating
350 Coating
HAZ
300
250 Base material
200
150
100
-3 -2 -1 0 1 2 3 4 5 6
Distance across the interface (mm)
Figure 2 - Hardness profile across the interface of C100 and C140 coatings.
Figure 3 shows the wear loss rate of coatings and gray cast iron tested at different
temperatures. It clearly shows that at room temperature the wear loss of coatings and gray cast
iron is very low, and almost independent of the substrate dilution. C100 coating displays a
little higher wear resistance than C140 coating, in agreement with its higher hardness. The
observed similar level of wear rate of the gray cast iron as compared to the Ni coatings at RT
is due to its higher amount of graphite that can act as solid lubricant, leading to a less volume
loss of material. It should be pointed out that the low hardness displayed by the gray cast iron
compared to the coated samples is compensated by the presence of graphite on the sliding
contact [13, 14].
At temperature of 550 ºC an abrupt increase of the wear rate of coatings was observed;
however, it dropped with further increase to 700 ºC. The first increase of the wear rate can be
attributed to the spontaneous oxidation, while, as it will be seen later, following decrease of
wear loss at 700 ºC could be related to the formation of high amount of oxides on the wear
track preventing further wear of the coatings. Besides, coating produced with high dilution
displayed much higher wear resistance at elevated temperatures than C100 coating. This
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behavior is a result of the low oxidation resistance of C140 coating that will contribute to the
formation of a high amount of oxides in the contact zone. A continuous increase of the wear
rate with increasing temperature was observed for the gray cast iron. This material showed
much higher wear rate than the Ni-based coatings, indicating that the latter should perform
better at high temperature. This is an expected result since rapid loss of the material
mechanical strength and increasing oxidation rate with temperature are typical of gray cast
iron.
450
RT
Wear rate (10 mm /N.m)
400 0
550 C
350 0
700 C
3
300
250
-6
200
150
100
50
0
C 100 C 140 Cast iron
Figure 3 - Variation of the wear rate of coatings and gray cast iron with test temperature.
The evolution of the friction coefficient of coatings and gray cast iron with increasing
test temperature is shown in Figure 4. As expected, all the friction curves displayed two
distinct regions: running-in stage and steady state stage. At RT, the running-in stage for both
Ni-based coatings is similar and characterized by strong increase of the friction coefficient to
the highest value of 0.74, followed by rapid decrease down to 0.41, after the first 1000 laps.
On the other hand, in gray cast iron, friction coefficient slightly increased over time reaching
the steady state after 4000 laps with a friction coefficient of about 0.39. Gray cast iron
exhibited relatively low friction coefficient at RT as compared to the Ni-based alloys.
According to the literature [13, 14], this behavior is related to the presence of graphite on the
worn surface that acts as a solid lubricant compensating the lower hardness displayed by the
material and thus leading to wear rate values similar to Ni-based coatings.
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A long running-in period was also noticed with a progressive increase of the friction
coefficient for the gray cast iron when the temperature was 550 ºC; however, friction was
much higher suggesting different contact conditions. Ni-based coatings displayed very short
running-in periods with friction stabilized at similar values as in RT tests. Finally, the tests at
700 ºC exhibited different trends between cast iron and coatings. In the first case, a small
decrease in the final friction coefficient value was achieved whereas an inverse situation was
observed for coated samples, with friction stabilized at values of approximately 0.6 and 0.65
for C100 and C140 coatings, respectively. Such trends suggest that different contact
conditions should have been taking place in both situations.
In many cases, especially at high temperatures, the friction curves displayed quite
unstable values. According to Kesavan and Kamaj [15], such behavior demonstrates the
presence of tribolayer on the sliding surfaces; the tribolayer is continually worn out and
formed again, which results in short-term friction coefficient oscillations.
0.6
0.5
0.4
0.3 RT
0
550 C
0.2 0
700 C
0.1
0 1000 2000 3000 4000 5000
a) Number of cycles
Figure 4 - Friction coefficient evolution with increasing test temperature of: a) C100 coating, b) C140 coating, c)
gray cast iron.
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Friction coefficient
0.6
0.5
0.4
0.3 RT
0
550 C
0.2 0
700 C
0.1
0 1000 2000 3000 4000 5000
b) Number of cycles
0.6
0.5
0.4 RT
0
550 C
0.3 0
700 C
0.2
0.1
0 1000 2000 3000 4000 5000
c) Number of cycles
Figure 4 (continued).
The changes of the wear rates and friction coefficients described above can be
attributed to a great number of factors, such as the chemical composition, physical properties
of materials, or applied test conditions. Therefore, the investigation of the interaction between
the specimen-counterpart pair was carried out. Scanning electron microscopy with energy
dispersive X-ray spectroscopy (SEM-EDS) and Raman spectroscopy were used to
characterize the dominant wear mechanisms and the wear debris originated by the tribological
testing.
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Typical SEM micrographs in secondary electrons of the worn surface of coatings and
gray cast iron tested against Al2O3 balls at RT are depicted in Figure 5. The wear tracks of Ni-
based coatings are similar; they present a rough morphology with evidence of tribo-oxidation.
The oxide adhered layer is cracked, which reveals its brittle nature. Raman spectra performed
at the oxidized worn surface of the coatings (see spectra a) and b) in Fig. 6 for C100 and C140
coatings, respectively) confirmed the presence of the oxides, namely Raman active modes
assigned to Ni-O [16] and C [17]; the latter almost inexistent in C100 sample spectrum as a
consequence of its much lower dilution. These results are in good agreement to EDS analysis
(see Fig. 5 d) for C100 coating as an example), which showed essentially strong signals of Ni
and O; vestiges of Fe, Si, Al and C were also detected. Similar wear debris was found on the
wear scars of the balls too. During the wear tests the friction, generated by the sliding of the
ball against the coating surface, causes the increase of the local temperature, mainly in the
protuberances contact, promoting the tribo-oxidation of the surface. During this phase, the
metal-ceramic contact will progressively decrease with more and more surface being covered
by an oxide phase. With the test running, the oxide tribo-layer will grow and the contact shall
be governed between this growing tribo-layer and the counterpart, until a steady state will
occur. Metal-ceramic contact gave rise to a higher friction coefficient that should decreases
until a constant and lower value corresponding to the same oxide, NiO as shown by Raman
spectroscopy, formed on the top surface. It explains the similar level of wear resistance and
friction values of the coatings even despite of their discrepancy in hardness and chemical
composition due to different dilution levels. The high hardness of the metallic coatings, in
conjunction with the protective effect of the hard NiO oxide, led to very low wear volumes.
Distinct wear behavior was found for the gray cast iron. Its wear track displays a very
clean surface with fine scratches parallel to the relative sliding movement (suggesting
abrasion wear mechanism) and some adherent wear debris. Further, in some places, black
regions can be noticed, which can be attributed to either the presence of graphite or its pull-
out from the material [18]. EDS analysis performed at the wear debris revealed that they
represent a mixture of original and oxidized material, including a weak signal from graphite
lamellas. Similar phases were detected on clean worn surface, although C peak was much
more intense there. Raman spectroscopy analysis showed the presence of hematite (α-Fe2O3)
[19] and nanocrystalline graphite through the characteristic D and G bands [20] (see spectrum
c) and d) of Fig. 6 for wear debris and clean surface, respectively). During sliding
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experiments the movement of the ball promotes the wear of surface specimen with production
of both oxide and graphite debris. Graphite will act as a solid lubricant in tribo-film,
promoting a decrease of friction coefficient and having, indirectly, an important impact on the
wear [13, 14]. As a consequence, similar level of wear resistance was observed between base
material and coatings, in spite of the lower hardness of gray cast iron.
d) O Kα
a)
Intensity (a.u.)
Ni Lα
1
Fe Lα
Si Kα
C Kα Al Kα
KV
0.4 0.8 1.2 1.6
b) 3.5
3
Depth (µm)
2.5
2
1.5
1
0.5
0
0.1 0.2 0.3 0.4 0.5
e) Scan length (µm)
c) 3.5
3
Depth (µm)
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NiO C Fe2O3
d)
Relative intensity
c)
b)
a)
200 400 600 800 1000 1200 1400 1600 1800 2000
-1
Raman shift [cm ]
Figure 6 - Typical Raman spectra performed on the worn surface of: a) C100 coating tested at RT, b) C140
coating tested at RT, c-d) adherent wear debris and clean surface of gray cast iron, respectively, tested at room
temperature.
When sliding occurred at 550 ºC, C100 and C140 coatings also showed oxides on the
contact surface. However, C140 wear track was almost fully covered whereas C100 exhibited
an irregular surface with a mixture of clean wear zones and oxidized ones with signs of
adhesive and abrasive wear (see Fig. 7 a-b), suggesting an incomplete formation of an oxide
transfer layer. The difference between the two samples can be interpreted considering their
oxidation behavior. In fact, C100 sample oxidizes through the formation of a protective
mixture of oxides based on B / Si over a Ni-based oxide layer [10]. Although these oxides are
protective in static oxidation conditions, B-oxide melts below 600 ºC [21, 22], which leads to
its easy removal hindering the formation of a stable oxide layer over the surface. As the
process requires some time, since temperatures are still low being oxidation rates reduced
[10], the destruction of the growing oxide layers is relatively easy and the formation of a
transfer layer is limited due to the dragging effect caused by the ball sliding and melted B-
oxide. Therefore, high wear volumes are expected and indeed the steep increase in the wear
rate can be noticed in Fig. 3. In the case of C140 sample, besides Ni-based oxides, iron oxide
is also formed during oxidation at these temperatures [10]. The oxidation rates are much
higher compared to C100 sample due to higher dilution. Therefore, a higher amount of oxide
debris can be produced facilitating the formation of a continuous oxide transfer layer. The
surface is thus protected quickly and, in spite of the much higher wear rate in comparison to
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RT test, the wear resistance is higher than that for C100 sample. Raman analysis performed at
the adhesive wear debris on C100 coating indicated that they were mainly Ni-O, although,
according to the EDS analysis, Fe-O, Si-O and Al-O should also coexist in much less amount.
It is in good agreement to our previous work on the oxidation behavior of this coating [10]. At
the clean zone (bottom part) of the worn surface EDS showed non-oxidized coating material.
For C140 coating, Raman analysis showed similar oxides but, in this case, a significant
amount of iron oxides (hematite and magnetite (Fe3O4)) were also detected, in good
agreement to higher level of dilution and to the oxidation studies where this oxide was found
to be the main constituent of the oxide scale [10]. Additionally, strong signals of C-based
phases were detected by EDS.
Gray cast iron was fully covered with an oxide layer at this temperature, with
scratches parallel to the movement of the sliding. The main oxides identified in the wear track
were hematite and magnetite (see Raman spectrum c) of Figure 8). The softening of the
material, induced by the temperature increase, together with the total depletion of graphite
from the track surface justify the sharp increase in the wear rate, as well as the increase in the
friction coefficient in comparison to results at RT.
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a) b)
Wear
debris
c) d)
e) f)
Figure 7 - SEM morphology in secondary and backscattered electrons of worn surface of: a-b) C100 coating, c-
d) 140 coating, e-f) cast iron; all tested at 550 ºC.
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Relative intensity
c)
b)
a)
200 400 600 800 1000 1200 1400 1600 1800 2000
-1
Raman shift [cm ]
Figure 8 - Raman spectra performed on the worn surface of: a) C100 coating tested at 550 ºC, b) C140 coating
tested at 550 ºC, c) gray cast iron tested at 550 ºC.
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a) b)
c) d)
e) f)
Figure 9 - SEM morphology in secondary and backscattered electrons of worn surface of: a-b) C100
coating, c-d) 140 coating, e-f) cast iron; all tested at 700 ºC.
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Relative intensity
c)
b)
a)
200 400 600 800 1000 1200 1400 1600 1800 2000
-1
Raman shift [cm ]
Figure 10 - Raman spectra performed on the worn surface of: a) C100 coating tested at 700 ºC, b) C140 coating
tested at 700 ºC, c) gray cast iron tested at 700 ºC.
Conclusions
Nickel based hardfaced coating deposited by plasma transferred arc on gray cast iron
with two dissimilar levels of dilution, were tribologically tested at room and high
temperatures and compared to uncoated gray iron. The results showed that the wear loss of
the coatings was independent of the substrate dilution at room temperature; tribo-oxidation
was the main wear mechanism. Despite its lower hardness, gray cast iron showed similar wear
rate as Ni-based coatings due to the lubricious effect of a C tribo-layer formed on the sliding
surfaces. Increase sliding test temperature to 550 ºC led to an abrupt increase of the wear rate
of the coatings and the gray cast iron due to the material softening and oxidation. Gray cast
iron showed extremely high wear values due to the continuous removal of the oxides, which
exposed the soft non-oxidized material. Ni-O based layer protected more efficiently the Ni-
based coatings. Further increase in test temperature to 700 ºC led to the improvement of the
wear resistance of coatings and an intensification of the wear on gray cast iron. The coating
produced with the highest dilution displayed the highest wear resistance at elevated
temperatures. This behavior was related to its lower oxidation resistance, which promoted the
easier formation of high amounts of oxide wear debris and their agglomeration into compact
oxide tribolayer acting as protection of wearing surfaces.
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Acknowledgments
This research is sponsored by FEDER funds through the program COMPETE – Programa
Operacional Factores de Competitividade – and by national funds through FCT – Fundação
para a Ciência e a Tecnologia –, under the projects PTDC/EME-TME/122116/2010, PEst-
C/EME/UI0285/2013 and CENTRO -07-0224 -FEDER -002001 (MT4MOBI), as well as the
grant (SFRH/BD/68740/2010).
References
[3] Vaßen R, Jarligo MO, Steinke T, Mack DE, Stöver D. Overview on advanced thermal
barrier coatings. Surf Coat Technol. 2010;205:938-42.
[4] Gatto A, Bassoli E, Fornari M. Plasma Transferred Arc deposition of powdered high
performances alloys: process parameters optimisation as a function of alloy and geometrical
configuration. Surf Coat Technol. 2004;187:265-71.
[5] Siva K, Murugan N, Logesh R. Optimization of weld bead geometry in plasma transferred
arc hardfaced austenitic stainless steel plates using genetic algorithm. Int J Adv Manuf
Technol. 2009;41:24-30.
[6] Guo C, Zhou J, Chen J, Zhao J, Yu Y, Zhou H. High temperature wear resistance of laser
cladding NiCrBSi and NiCrBSi/WC-Ni composite coatings. Wear. 2011;270:492-8.
[7] Gurumoorthy K, Kamaraj M, Rao KP, Rao AS, Venugopal S. Microstructural aspects of
plasma transferred arc surfaced Ni-based hardfacing alloy. Materials Science and
Engineering: A. 2007;456:11-9.
142
Filipe Daniel Fernandes
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[15] Kesavan D, Kamaraj M. The microstructure and high temperature wear performance of a
nickel base hardfaced coating. Surf Coat Technol. 2010;204:4034-43.
[17] Jian S-R, Chen Y-T, Wang C-F, Wen H-C, Chiu W-M, Yang C-S. The Influences of H2
Plasma Pretreatment on the Growth of Vertically Aligned Carbon Nanotubes by Microwave
Plasma Chemical Vapor Deposition. Nanoscale Research Letters. 2008;3:230-5.
143
Filipe Daniel Fernandes
Annex E
[18] Pandya SN, Nath SK, Chaudhary GP. Friction and Wear Characteristics of TIG
Processed Surface Modified Gray Cast Iron2009.
[19] de Faria DLA, Venâncio Silva S, de Oliveira MT. Raman microspectroscopy of some
iron oxides and oxyhydroxides. Journal of Raman Spectroscopy. 1997;28:873-8.
[21] Li YQ, Qiu T. Oxidation behavior of boron carbide powder. Materials Science and
Engineering: A. 2007;444:184-91.
[22] Lavrenko VA, Gogotsi YG. Influence of oxidation on the composition and structure of
the surface layer of hot-pressed boron carbide. Oxid Met. 1988;29:193-202.
144
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Annex F
Annex F
F. Fernandes, A. Ramalho, A. Loureiro, J.M. Guilemany, M. Torrell, A.
Cavaleiro, Influence of nanostructured ZrO2 additions on the wear resistance of Ni-
based alloy coatings deposited by APS process, Wear, 303 (2013) 591-601.
145
Filipe Daniel Fernandes
Wear 303 (2013) 591–601
Wear
journal homepage: www.elsevier.com/locate/wear
art ic l e i nf o a b s t r a c t
Article history: In the present investigation, the influence of the addition of nanostructure zirconia particles on the
Received 19 December 2012 microstructure, micro-hardness and wear performance of a Ni-based alloy (Colmonoy 88) deposited by
Received in revised form atmospheric plasma spraying (APS) on low carbon steel has been reported. Two different procedures
8 April 2013
were tested: (i) spraying powders of Colmonoy 88 and zirconia mixed by mechanical alloying and (ii)
Accepted 15 April 2013
spraying powders separately using a dual powder injection system available at the APS equipment. The
Available online 22 April 2013
microstructure and the mechanical properties of coatings were characterized by scanning electron
Keywords: microscopy/energy dispersive X-ray analysis (SEM-EDS), X-ray diffraction (XRD) and micro-hardness
Sliding wear measurements. The tribological properties were evaluated at room temperature in reciprocating sliding
Thermal spray coatings
wear equipment. The results indicate that the as-sprayed modified coatings were mainly composed by
Metal–matrix composite
Ni, Ni–Cr–Fe, Cr23C6, Cr5B3, and tetragonal zirconia. Evenly distribution of zirconia can be seen in the
Surface analysis
Wear testing coatings produced by powders prepared by mechanical alloying, while dispersive ones can be seen in the
other case. Hardness and wear resistance of coatings is increased with nanostructured zirconia additions,
while their friction coefficient is decreased. Coatings produced with mechanical alloying show the
highest wear resistance of all tested coatings. Nanostructured ZrO2 coating displays the worst wear
resistance.
& 2013 Elsevier B.V. All rights reserved.
1. Introduction behavior of cobalt, iron and nickel alloys coatings have been
studied using several alloy compositions, several processes of
Improvement of the surface properties by thermal spraying deposition and different substrates. Regarding the wear resistance
hard and wear resistant materials is a commonly used industrial of these coatings, some special composite systems were studied
practice [1–3]. The base material provides the overall mechanical and developed. The incorporation of hard and stable carbides and
strength of the components while coatings provide a way of oxide phases, such as CrC, SiC, WC, Al2O3, TiO2, ZrO2, CeO2, etc, in
extending the limits of their use at the upper end of their these coatings, has been reported as a way to improve their wear
performance capabilities. resistance. For example, Hou et al. [15] showed the beneficial
A wide variety of coating materials such as cobalt, iron and effect of nano-Al2O3 particles on the microstructure and wear
nickel alloys can be successfully applied by thermal spraying resistance of a nickel-based alloy coating deposited by plasma
technologies to protect the surface of components subjected to transferred arc (PTA). Harsha et al. [1] studied the influence of CrC
harsh environments [4–6]. Superalloys and cermet APS (atmo- on the microstructure, microhardness and wear resistance of a
spheric plasma spraying) coatings have been widely employed to nickel-based alloy deposited by flame spraying process. They
improve the oxidation, corrosion, abrasion and wear resistance of observed that the wear resistance of the Ni modified coating is
engineering components, such as plungers, molds, wearing plats, increased relatively to the CrC-free Ni-based alloy. Wang et al. [16]
turbines, tools, etc, whose surface is submitted to extreme tribo- investigated the effect of three nano-particles additions (Al2O3,
logical conditions in service [7–10]. Nickel-based alloys have been SiC, CeO2) on the high temperature wear behavior of a Ni-based
especially used in the protection of parts due to their unique alloy coatings produced by laser cladding technique and reported
combination of properties (mechanical, tribological, and high that the addition of these nanoparticles increased the wear
temperature properties) [11–14]. The microstructure and the wear resistance of the coatings. Regarding the effect of incorporation
of ZrO2 particles in a Ni matrix, it has also been extensively studied
[17–20]; however, only electrodeposited coatings have been con-
n
Corresponding author. Tel.: +351 239 790 745; fax: +351 239 790 701. cerned. Despite of these studies to improve the wear resistance of
E-mail address: filipe.fernandes@dem.uc.pt (F. Fernandes). Ni electrodeposited coatings, at our knowlodgment no reports
0043-1648/$ - see front matter & 2013 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.wear.2013.04.012
592 F. Fernandes et al. / Wear 303 (2013) 591–601
have been published as regards to the effect of nanostructured hardness, wear and oxidation resistance, its use as reinforcement
zirconia additions on the properties of Ni-based thermal sprayed material in nickel-based alloy coatings deposited by thermal
coatings. Since zirconia is a ceramic material of high level of spraying processes should be investigated.
Therefore, the main goal of the present work was to study the
Table 1 effect of nanostructured zirconia additions on the microstructure,
Nominal chemical composition (wt.%) of the as-received powders. micro-hardness and wear performance of a nickel-based hardfacing
alloy deposited on low carbon steel by atmospheric plasma spraying
Powders W B C Cr Fe Si Ni
(APS). The coatings with nanostructured zirconia additions were
Colmonoy 88 13.5 2.6 0.6 14.1 3.5 3.7 Balance produced by spraying either powders prepared by mechanical alloy-
ing or separate powders using a dual system of powder injection.
Y2O3
The microstructure, the mechanical and tribological properties
HfO2 ZrO2
were characterized by scanning electron microscopy/energy disper-
sive X-ray analysis (SEM-EDAX), X-ray diffraction (XRD), Vickers
Nanostructured ZrO2 2.5 7.5 Balance
indentation tests and reciprocating sliding wear tests. All the
Table 2
Plasma spraying parameters of the dissimilar coatings.
Fig. 1. SEM morphologies of: (a) nickel-based alloy powders, (b) nanostructured zirconia powders, (c) powders prepared by mechanical alloying, (condition with 40% ZrO2).
F. Fernandes et al. / Wear 303 (2013) 591–601 593
properties of the doped coatings were compared with the un- in the MA method the substrates were placed in a rotatory table and
modified nickel and nanostructured zirconia coatings. the spraying was performed by combining this rotation with the
vertical movement of the torch. Before final depositions, preliminary
coatings were produced changing the main deposition parameters
2. Experimental procedure and, then, characterized in order to properly select the depositions
conditions that allowed achieving the best results. Moreover, during
2.1. Materials depositions the substrates were cooled down by two pressurized air
guns keeping the temperature below 150 1C, in order to avoid cracking
In this investigation commercial powders of a Ni-based alloy and residual stresses formation.
(Colmonoy 88) and nanostructured zirconia from the “Wall Colmonoy” Powders mixtures were prepared by mechanical alloying in a
and “Innovnano” companies, respectively, were used as feedstock to Fritsch planetary ball mill using a 250 ml hardened steel vial and
produce coatings with Atmospheric Plasma Spraying process (APS) fifteen balls with 20 mm diameter of the same material. Two propor-
onto a low carbon steel substrate (AISI 8620). The Ni-based alloy tions of nanostructured zirconia were added to the nickel-based alloy
powders have been characterized as a spherical morphology with an powder, 20% and 40% in volume, i.e. 5.1 and 12.5 gr of nanostructured
average size D50 of 45 mm and density of 7900 (kg/m3), whilst ZrO2 powder for 94.9 and 87.4 gr of nickel-based alloy powder,
nanostructured ZrO2 powders have an average primary particle size respectively. The MA process was carried out at 400 rpm for 2 h in a
of 60 nm with granule size D50 of 60 mm and density 1700 (kg/m3). protective atmosphere of argon. After 1 h of milling, the process was
The chemical compositions of the as-received powders are displayed interrupted for 10 min to cool the vial and to reserve rotation. Finally,
in Table 1. The plasma sprayed coatings were deposited with an APS after milling, the powder was mechanical sieved to obtain particles
A3000 system from “Sulzer Metco”, using either pure as-received with a size distribution in the range [28–71 mm]. Two different
powders and a dual system for powder injection available at the coatings also containing nanostructured zirconia additions of 20%
APS equipment (designated as “Dual”) or mixtures of powders and 40% were also performed by using the dual system of powder
prepared by mechanical alloying (MA) process (designated as “MA”). injection available at the APS equipment. In this case the proportion
The substrate material was grit-blasted with alumina grade 24 from was easily achieved by controlling the flow of powders through the
“Alodur Germany” before coating (which induced a corresponding rotation control of the powder feeders. Henceforth, coatings produced
superficial roughness (Ra) of 5 mm), to increase the surface roughness using as received pure powder of nickel-based alloy and nanostruc-
and achieve the proper mechanical interlocking between the coating tured zirconia will be designated as “Ni”, and “ZrO2”, respectively,
and the substrate. For pure colmonoy and zirconia coatings and Dual coatings produced with powders prepared by mechanical alloying
procedure, the substrates were fixed positioned, being the spraying of using 20% and 40% of nanostructured ZrO2 will be termed as “Ni+20%
the surface ensured by the movement of the torch. On the other hand, MA” and “Ni+40% MA”, respectively, and coatings deposited using the
dual system of powder injection with 20% and 40% of nanostructured
ZrO2 will be designated as “Ni+20% Dual” and “Ni+40% Dual”,
respectively. The process parameters employed in the dissimilar
coating depositions are shown in Table 2.
Fig. 3. Energy dispersive spectroscopy analysis (EDS) of a powder produced by mechanical alloying (condition Ni+40% ZrO2).
594 F. Fernandes et al. / Wear 303 (2013) 591–601
2.3. Wear tests of signal used was 60 s. A soda-lime glass sphere with 500 HV0.5 of
hardness and 10 mm in diameter was used as counterpart, in order
Wear behavior of APS coatings was studied using reciprocating to study the interaction of glass with the different coatings. After
wear equipment. An harmonic wave generated by an eccentric and wear tests, the volume loss was estimated through the transverse
rod mechanism imposed a stroke length of 2.05 mm at a frequency profile of each wear scar at the middle and near the end of both sides
of 1 Hz. A detailed description of the equipment can be found in of the scar. To ensure the reproducibility of results, a set of three scars
reference [21]. Before wear tests, the surface of coatings were for each test condition was performed at the coating, and the volume
grounded and then polished using a 3 mm diamond past. A superficial loss calculated by using the volume average of these marks. An
roughness Ra of: 0.51, 0.65, 0.49 and 0.98 mm was measured approach of the Archard's law was applied; see Eq. (1), in order to
respectively for Ni, Ni Dual, Ni MA and ZrO2 coatings. In the scope obtain the specific wear rate. In this equation, V represents the wear
of the present study, the values of normal load applied to the coated volume, P is the normal load, k is the specific wear rate and l the
samples were 5, 7, 8.5 and 10 N. All the tests were conducted at a sliding distance. To estimate the error of the measurements, a
constant rotational speed of 190 rpm during 2 h. The acquisition time statistical analysis described elsewhere by A. Ramalho [22] was used.
Fig. 4. SEM morphology of the: (a) Ni-based alloy coating, (b) nanostructured ZrO2 coating, (c) and (d) Ni+20% Dual and Ni+40% Dual, respectively, (e) and (f) Ni+20% MA
and Ni+40% MA.
F. Fernandes et al. / Wear 303 (2013) 591–601 595
After wear tests, the morphology of depressions was observed and Ni and ZrO2 powders clearly display spherical-shaped particles, while
characterized by scanning electron microscopy. MA powders have irregular shapes. XRD patterns in Fig. 2 shows that
V ¼kPl ð1Þ Ni powders consist essentially in a Ni solid solution (ICDD card 01-
1258) with some traces of chromium carbide (ICDD card 85-1281)
and tungsten (ICDD card 47-1319). This indexation is in agreement
3. Results and discussion with the results from the literature where similar phases were
detected [23]. Nanostructured ZrO2 powders are essentially formed
3.1. Powders characterization by tetragonal zirconia. In respect to MA powders XRD pattern reveal
the same phases indexed from the pure Ni and nanostructured ZrO2
Fig. 1(a–c) shows SEM morphologies of nickel-based alloy, powders, suggesting only physical mixing, not having occurred
nanostructured zirconia and one of the MA prepared powders. chemical reactions. Fig. 3 shows a SEM-EDS analysis performed at a
Fig. 5. SEM-EDS analysis performed at the: (a) Ni coating, (b) Ni+40% Dual coating, (c) Ni+40% MA coating. SEM-EDS spectra of: (d) point 1, (e) zone 2, (f) point 3, (g) zone 4.
596 F. Fernandes et al. / Wear 303 (2013) 591–601
Fig. 8. Volume loss of: (a) coatings, (b) ball; after wear tests as function of the
applied load.
Fig. 7. Microhardness of coatings.
3.2. Microstructure
of substrate, then deform, solidify and transform into lamellae. the deposition. In fact, according to the literature the nature of
During spraying, the particles can be in all of the following powders has a detrimental effect in their interaction with the
states on impact: fully molten, superheated, semi-molten and plasma spraying flame and, therefore, on the properties of the
molten then re-solidified [24,25], depending on the deposition coatings [28]. The low level of melting particles can also explain
parameters. It is well known that in plasma spraying the electric the high level of porosity observed in these coatings.
arc current, the primary plasma gas flow rate, the second plasma For better characterization and perception of the different
gas flow rate and powder size are the main parameters that coatings microstructure, magnifications of the Ni, Ni+40% Dual
influence the in-flight particle behaviors. From Table 2, comparing and Ni+40% MA coatings, as well as EDS analyses are shown in
the process parameters for the Ni and Ni+20% and 40% ZrO2 Dual Fig. 5. The EDS analysis revealed that Ni coating is essentially
coatings, differences can be observed in relation to the spray formed by Ni, W, Cr and Fe, which matches well with the chemical
distance, primary and secondary gas flow rates, and powder feed composition of the as received Ni powders. Moreover, small
rate. According to the literature, an increase of argon flow rate particles rich in W (white phase) can be observed evenly dis-
leads to a gentle decrease of particle temperature but a rapid tributed in the microstructure. The semi-melted particles in the
increase of velocity [26]. On the other hand, increasing the microstructure of the Ni+40% Dual coating are from the same
hydrogen flow rate increases both, temperature and velocity, nature of the Ni coating matrix, therefore, they can be identified as
though temperature is more sensitive to the hydrogen flow rate semi-melted Colmonoy powders. Similar to Ni coating a white
than velocity. So, due to the lower primary gas flow rate (Ar) and phase rich in W is also detected in the microstructure, co-existing
higher secondary gas flow rate (H2) used in the Ni +ZrO2 Dual with a coarse gray phase, entrapped between the boundaries of
coatings, would be expected a great level of melting particles and the semi-melted particles, revealed by EDS to be zirconia. Such a
consequently an homogenous structure, if the particles do not distribution is in the basis of a brittle behavior. Contrary to this
oxidize during the in-flight time. However, such fact was not coating, Ni+40% MA displays a compact and homogeneous micro-
observed. The lower primary gas flow rate (Ar) used in the structure, with small zirconia particles evenly distributed along
production of these coatings might suggest that independently the microstructure improving the coating toughness. Therefore,
of the high melting state of particles, the velocity given by the flow from these results, it can be inferred that the addition of nano-
could be not enough to accelerate the particles onto the specimen structured zirconia can be successfully achieved by using both
surface. Nevertheless, the same process parameters were used in procedures, powders prepared by mechanical alloying or dual
the deposition of the Ni+20% MA coating, with exception of the projection. However, in coatings deposited with MA powders
spraying distance, and the microstructure revealed to be homo- zirconia is finer and more homogeneously distributed in the
genous and compact. In fact, the spraying distance also plays an Ni-matrix giving rise to better mechanical performance.
important role in the thermal treatment of particles during their XRD patterns of the coatings are displayed in Fig. 6. As the
in-flight time [27]. Thereby, for short spraying distances, due to spectrum reveals, the nickel-based coating consists mainly of Ni,
the short exposure of particles in the flame, the energy added by Ni-Cr-Fe, Cr23C6 and Cr5B3 phases. These results are in accordance
the flame could be not enough to semi-molten or molten the with the literature where similar phases were identified [23].
particles, giving rise to a microstructure similar to those observed Moreover, all the coatings with nanostructured zirconia additions
for these coatings. Furthermore, it can also be speculated that the showed diffraction lines corresponding to the same phases identi-
two different ways to add zirconia at the nickel alloy can change fied at both the Ni and ZrO2 coatings. In the latter, is mainly
the entropy of the system, and therefore change the running-on of formed by tetragonal zirconia.
Fig. 10. Variation of the friction force of coatings with applied load for: (a) Ni coating, (b) Ni+40% Dual coating, (c) Ni+40% MA coating (d) nanostructured ZrO2 coating.
598 F. Fernandes et al. / Wear 303 (2013) 591–601
3.3. Micro-hardness domains in a soft matrix gives rise to tougher materials, which
combined lead to a decrease of the volume loss due to wear. Similar
Fig. 7 shows that nanostructured ZrO2 additions increases the behavior was observed by other author after adding hard particles to a
hardness of the nickel-based alloy. The same behavior was reported nickel based alloy [1,15]. In the case of Ni+ZrO2 Dual coatings
by other authors after addition of hard particles such as (Al2O3, CeO2, nanostructured hard ZrO2 phases have initially a similar influence,
SiC, WC) at a nickel alloy [1,3,16]. However, MA coatings have higher by increasing the hardness. However, the agglomeration of ZrO2
hardness than Dual ones. This result is coherent with either the particles, localized in-between Ni-based lamellae makes microstruc-
lower porosity or the finer and more homogeneous distribution of ture less tough with its consequent detrimental effect on the wear
nanostructured zirconia in Ni+ZrO2 MA coatings. performance when high contents of the hard phase exist.
In order to normalize the wear results, they were plotted in Fig. 9
3.4. Wear behavior in terms of material volume loss as function of the product of sliding
distance by the applied load. For each set of data, a straight linear
The volume loss due to wear of the coatings and counterpart trend was fitted in order to obtain the specific wear rate from the
increases with increasing applied load, as it is shown in Fig. 8(a) and slope (see Eq. 1). Further, a complete reliability analysis of the specific
(b), respectively, being this trend more notorious for ZrO2 coating. As it wear rate is presented in Table 3 for a confidence interval of 90%.
will be shown later, the higher volume loss for this coating can be Because of the high volume loss displayed by the ZrO2 coating, only
correlated with the brittle behavior as displayed on its worn surface. the point produced with lower load and the beginning of the straight
Nanostructured zirconia additions have a positive effect, decreasing adjustment were plotted in Fig. 9. Crossing the information of Fig. 9
the volume loss of material, independently of the deposition proce- and Table 3 it is possible conclude the specific wear rate of
dure used. However, efficiency seems to be higher in Ni+ZrO2 MA Ni-modified coatings is lower than that of the un-modified Ni coating.
coatings which display higher decreases of the volume loss than Ni Ni+40% ZrO2 dual coating display the lowest specific wear rate, being
+ZrO2 Dual coatings. Furthermore, increasing zirconia content in MA this result attributed to the higher hardness displayed by this coating.
coatings continues to have a beneficial effect whereas the inverse is Further, coatings produced by the dual system of powder injection
observed in Dual coatings. The better wear performance of Ni+ZrO2 display higher specific wear rate than coatings produced by powders
MA coatings can be attributed to the combination of a higher prepared by mechanical alloying. ZrO2 coating displays the highest
hardness, an evenly distribution of smaller zirconia domains in the specific wear rate, being one and two orders of magnitude greater
Ni-based matrix and a more compact microstructure. In fact, the than the un-modified and modified coatings, respectively.
presence of hard phases increases the overall hardness of the coating The average values of the frictional force of the Ni, Ni+40% Dual,
and, on the other hand, smaller and uniformly distributed hard Ni+40% MA and ZrO2 coatings linearly increases with the applied
Fig. 11. SEM morphologies of the worn surfaces tested at 8.5 N load of the: (a) Ni coating, (b) Ni+40% dual, (c) Ni+40% MA, (d) ZrO2 coating.
F. Fernandes et al. / Wear 303 (2013) 591–601 599
load, as shown in Fig. 10. This is in good agreement with Amontons– with the nanostructured zirconia additions, independently of the
Coulomb model. By fitting a linear trend to the experimental deposition method. Pure nanostructured zirconia shows higher
results, the friction coefficient calculated from the slope decreases friction coefficient than composite coatings, but lower than the
Fig. 12. Magnification of the worn surfaces tested with a load of 8.5 N of the: (a) Ni coating, (b) Ni+40% dual, (c) Ni+40% MA, (d) ZrO2 coating. SEM-EDS spectra of:
(e) point A, (f) point B.
600 F. Fernandes et al. / Wear 303 (2013) 591–601
pure nickel one. The changes in the specific wear rate and friction Dual coatings show a mix morphology of pure Ni and ZrO2
coefficient described above are only partially correlated with the coatings, i.e. adhered debris can be observed over a rough
applied test conditions and properties of the materials, which can brittle-appearance surface, whose number decreases with nanos-
produce dissimilar wear mechanisms during wear tests, and there- tructured zirconia additions. In these coatings, zirconia particles
fore dissimilar volume loss of coatings. These reasons have encour- are localized in the boundaries of the semi-melted Ni-based
aged the investigation of the interaction between the pair powders (see Fig. 4c and d). Despite of the particles operating as
specimens-counterpart. Scanning electron microscopy (SEM) ana- reinforcement of the coatings, leading to a lower volume loss, the
lysis was conducted on the dissimilar worn surfaces of coatings in area of the nickel exposed to the glass sphere remains higher.
order to identify the main wear mechanisms produced. Therefore, it is expected similar debris adhesion as for Ni coating.
Typical SEM morphologies of worn surfaces of the coatings At the same time, due to the continuous removal of material, ZrO2
tested with 8.5 N against a glass sphere are shown in Fig. 11. particles in the boundaries of the semi-melted powders loses their
Distinct wear mechanisms were identified in the worn surfaces. In support capacity being removed from the boundaries. These
Ni and Ni+ZrO2 Dual coatings significant amount of adherent particles will slide through the wear track acting as a removal of
material is shown in the wear scar, while Ni+ZrO2 MA and ZrO2 the wear debris. Increasing the volume of nanostructured zirconia
coatings show clean worn surfaces. EDS analysis performed in from 20 to 40%, higher levels of brittleness in the boundaries are
adhered debris, marked in Fig. 12(a) and (b) by letters A and B, achieved, increasing the volume loss of material, as reported
allows concluding that they are made of silicon oxides and Ni, before, and giving rise to the rough and brittle appearance shown
suggesting a mechanical mixing by the friction process, joining the below the adhered debris. The higher level of porosity of these
softer Ni phase with the silica type wear debris from the ball wear. coatings also accounts for higher levels of volume loss of material.
Being softer, the compact wear debris can stick at the worn track The coatings produced from MA powders show also clean wear
as well as at the ball wear scar, as can be observed in Fig. 13. surfaces but with a different failure type from pure zirconia coating. Ni
During contact the friction generated by the movement of the +ZrO2 MA coatings reveal surfaces with scratches identifying some
counterpart against the surface specimen, causes an increase of sort of grooving abrasion, see Fig. 12(c). The uniform distribution of
temperature promoting local plastic deformation. At this time, small nanostructured zirconia particles and the higher hardness and
both materials will adhere to each other. In the wear track of the toughness of these coatings avoids the plastic deformation, the
Ni coating, see Fig. 12(a), the wear debris tend to be incorporated adhesion and the liberation of large wear debris either from the
in the specimens, more precisely in the boundaries of the semi- coating or the counterpart. The smooth worn surfaces can be in the
melted powders, remaining adherent and leading to removal of basis of the low friction coefficient [30]. In these coatings, the matrix
material. This mechanism explains simultaneously the high fric- cutting occurs through a process of two body abrasion, as shown
tion values reported for Ni coating. On the other hand, ZrO2 Fig. 12(c). Despite of the dissimilar wear mechanisms observed in the
coating shows clean wear scars without traces of adhered material. wear track of the coatings, during the wear tests it was observed that
The surface has a brittle appearance, where micro cracks can be most of the material resulting from the wear was blown out, due to
seen, as shown in Fig. 12(d), revealing a rough surface from which the movement of the ball.
particles of several tens of micrometers were detached. This
morphology suggests mechanical interlocking during the contact
which can explain the higher values of the friction coefficient.
Furthermore, it should promote a strong abrasion in the counter- 4. Conclusions
part, giving rise to much higher wear rate as shown in Fig. 8(b).
Moreover, the evolution of the wear amount with the increase of The influence of nanostructured zirconia additions on the
normal load, (see Fig. 8), allow identifying a made severe wear microstructure, micro-hardness and wear performance of a
volume for loads higher than 7 N, which should be correlated to a nickel-based hardfacing alloy deposited by atmospheric
transition for a made brittle behavior. Similar behavior was plasma spraying (APS) on low carbon steel is reported in the
observed by Chen et al. [29] during the evaluation of present paper.
un-lubricated wear properties of plasma-sprayed nanostructures The coatings containing nanostructured zirconia additions
and conventional zirconia coatings. The wear scars of Ni+ZrO2 were produced either using powders prepared by mechanical
alloying or using separate powders that are sprayed simulta-
neously via a dual system of powder injection.
Nanostructured zirconia can be successfully incorporated into
nickel-based alloy powders by mechanical alloying. Coatings
produced with these powders present evenly distributed zir-
conia as opposed to coatings produced by separate powder
injection, where the nanostructured zirconia particles were
entrapped in the boundaries of the semi-melted Ni lamellae.
The microstructure of the nickel coating consists mainly of Ni,
Ni–Cr–Fe, Cr23C6 and Cr5B3 phases. The same diffraction lines
were identified in the modified nickel coatings with tetragonal
zirconia, independently of the procedure used in depositions.
The hardness and wear behavior of coatings was improved
with nanostructured zirconia additions while the friction
coefficient is decreased.
The nanostructured ZrO2 coating shows the highest hardness
values but also the highest specific wear rate, one to two orders
of magnitude higher than the other prepared coatings, in
particular those deposited with MA powders.
Fig. 13. Optical micrograph of the ball wear scars tested against the Ni coating with Adhesive wear mechanism was observed in the worn surface of
7 N of load. the nickel coating and coatings prepared by Dual deposition,
F. Fernandes et al. / Wear 303 (2013) 591–601 601
while abrasive wear was observed in ZrO2 coating and coatings [13] C. Guo, J. Zhou, J. Chen, J. Zhao, Y. Yu, H. Zhou, High temperature wear
deposited using MA powders. resistance of laser cladding NiCrBSi and NiCrBSi/WC–Ni composite coatings,
Wear 270 (2011) 492–498.
[14] F. Fernandes, A. Cavaleiro, A. Loureiro, Oxidation behavior of Ni-based coatings
deposited by PTA on gray cast iron, Surface and Coatings Technology 207
Acknowledgments (2012) 196–203.
[15] Q.Y. Hou, Z. Huang, J.T. Wang, Influence of nano-Al2O3 particles on the micro-
The authors wish to express their sincere thanks to the structure and wear resistance of the nickel-based alloy coating deposited by
plasma transferred arc overlay welding, Surface and Coatings Technology 205
Portuguese Foundation for the Science and Technology (FCT),
(2011) 2806–2812.
through COMPETE program from QREN and to FEDER, for financial [16] H.-y. Wang, D.-w. Zuo, M.-d. Wang, G.-f. Sun, H. Miao, Y.-l. Sun, High temperature
support in the aim of the Project number “13545”, as well as for the frictional wear behaviors of nano-particle reinforced NiCoCrAlY cladded coatings,
Grant (SFRH/BD/68740/2010). Transactions of Nonferrous Metals Society of China 21 (2011) 1322–1328.
[17] W. Wang, H.T. Guo, J.P. Gao, X.H. Dong, Q.X. Qin, XPS, UPS and ESR studies on
the interfacial interaction in Ni–ZrO2 composite plating, Journal of Materials
References Science 35 (2000) 1495–1499.
[18] K.F. Zhang, S. Ding, G.F. Wang, Different superplastic deformation behavior of
nanocrystalline Ni and ZrO2/Ni nanocomposite, Materials Letters 62 (2008)
[1] S. Harsha, D. Dwivedi, A. Agarwal, Influence of CrC addition in Ni–Cr–Si–B
719–722.
flame sprayed coatings on microstructure, microhardness and wear behaviour,
[19] F.Y. Hou, W. Wang, H.T. Guo, Effect of the dispersibility of ZrO2 nanoparticles in
International Journal of Advanced Manufacturing Technology 38 (2008)
93–101. Ni–ZrO2 electroplated nanocomposite coatings on the mechanical properties
[2] J.M. Miguel, J.M. Guilemany, S. Vizcaino, Tribological study of NiCrBSi coating of nanocomposite coatings, Applied Surface Science 252 (2006) 3812–3817.
obtained by different processes, Tribology International 36 (2003) 181–187. [20] W. Wang, F.Y. Hou, H. Wang, H.T. Guo, Fabrication and characterization
[3] J.C. Miranda, A. Ramalho, Abrasion resistance of thermal sprayed composite of Ni–ZrO2 composite nano-coatings by pulse electrodeposition, Scripta
coatings with a nickel alloy matrix and a WC hard phase. Effect of deposition Materialia 53 (2005) 613–618.
technique and re-melting, Tribology Letters 11 (2001) 37–48. [21] A. Ramalho, P.V. Antunes, Reciprocating wear test of dental composites against
[4] A.K. Basak, S. Achanta, J.P. Celis, M. Vardavoulias, P. Matteazzi, Structure and human teeth and glass, Wear 263 (2007) 1095–1104.
mechanical properties of plasma sprayed nanostructured alumina and [22] A. Ramalho, A reliability model for friction and wear experimental data, Wear
FeCuAl–alumina cermet coatings, Surface and Coatings Technology 202 269 (2010) 213–223.
(2008) 2368–2373. [23] J.G. La Barbera-Sosa, Y.Y. Santana, E. Moreno, N. Cuadrado, J. Caro, P.O. Renault,
[5] Y. Wu, S. Hong, J. Zhang, Z. He, W. Guo, Q. Wang, G. Li, Microstructure and E. Le Bourhis, M.H. Staia, E.S. Puchi-Cabrera, Effect of spraying distance on the
cavitation erosion behavior of WC–Co–Cr coating on 1Cr18Ni9Ti stainless steel microstructure and mechanical properties of a Colmonoy 88 alloy deposited
by HVOF thermal spraying, International Journal of Refractory Metals and by HVOF thermal spraying, Surface and Coatings Technology 205 (2010)
Hard Materials 32 (2012) 21–26. 1799–1806.
[6] R.A. Mahesh, R. Jayaganthan, S. Prakash, Oxidation behavior of HVOF sprayed [24] X.B. Zhao and Z.H. Ye, Microstructure and wear resistance of molybdenum
Ni–5Al coatings deposited on Ni- and Fe-based superalloys under cyclic based amorphous nanocrystalline alloy coating fabricated by atmospheric
condition, Materials Science and Engineering: A 475 (2008) 327–335. plasma spraying, Surface and Coatings Technology http://dx.doi.org/10.1016/j.
[7] H. Liao, B. Normand, C. Coddet, Influence of coating microstructure on the surfcoat.2012.05.127, in press.
abrasive wear resistance of WC/Co cermet coatings, Surface and Coatings [25] T. Sidhu, S. Prakash, R.D. Agrawal, Studies on the properties of high-velocity
Technology 124 (2000) 235–242. oxy-fuel thermal spray coatings for gigher temperature applications, Materials
[8] C.-J. Li, C.-X. Li, Y.-Z. Xing, M. Gao, G.-J. Yang, Influence of YSZ electrolyte
Science 41 (2005) 805–823.
thickness on the characteristics of plasma-sprayed cermet supported tubular
[26] J.C. Fang, H.P. Zeng, W.J. Xu, Z.Y. Zhao, L. Wang, Prediction of in-flight particle
SOFC, Solid State Ionics 177 (2006) 2065–2069.
behaviors in plasma spraying, World Academy of Materials and Manufacturing
[9] M.H. Staia, T. Valente, C. Bartuli, D.B. Lewis, C.P. Constable, A. Roman, J. Lesage,
Engineering 18 (2006) 283–286.
D. Chicot, G. Mesmacque, Part II: tribological performance of Cr3C2–25% NiCr
[27] O. Sarikaya, Effect of some parameters on microstructure and hardness of
reactive plasma sprayed coatings deposited at different pressures, Surface and
Coatings Technology 146–147 (2001) 563–570. alumina coatings prepared by the air plasma spraying process, Surface and
[10] G.J. Gibbons, R.G. Hansell, Thermal-sprayed coatings on aluminium for mould Coatings Technology 190 (2005) 388–393.
tool protection and upgrade, Journal of Materials Processing Technology 204 [28] P. Fauchais, Understanding plasma spraying, Journal of Physics D: Applied
(2008) 184–191. Physics 37 (2004) R86–R108.
[11] J.H. Chang, C.P. Chang, J.M. Chou, R.I. Hsieh, J.L. Lee, Microstructure and bonding [29] H. Chen, S. Lee, X. Zheng, C. Ding, Evaluation of unlubricated wear properties
behavior on the interface of an induction-melted Ni-based alloy coating and AISI of plasma-sprayed nanostructured and conventional zirconia coatings by SRV
4140 steel substrate, Surface and Coatings Technology 204 (2010) 3173–3181. tester, Wear 260 (2006) 1053–1060.
[12] W. Li, Y. Li, C. Sun, Z. Hu, T. Liang, W. Lai, Microstructural characteristics and [30] L.M. Chang, M.Z. An, H.F. Guo, S.Y. Shi, Microstructure and properties of Ni–Co/
degradation mechanism of the NiCrAlY/CrN/DSM11 system during thermal nano-Al2O3 composite coatings by pulse reversal current electrodeposition,
exposure at 1100 1C, Journal of Alloys and Compounds 506 (2010) 77–84. Applied Surface Science 253 (2006) 2132–2137.
Annex G
Annex G
F. Fernandes, A. Loureiro, T. Polcar, A. Cavaleiro, "The effect of increasing V
content on the structure, mechanical properties and oxidation resistance of Ti-Si-V-
N films deposited by DC reactive magnetron sputtering, Applied Surface Science,
289, (2014) 114-123.
159
Filipe Daniel Fernandes
Applied Surface Science 289 (2014) 114–123
a r t i c l e i n f o a b s t r a c t
Article history: In the last years, vanadium rich films have been introduced as possible candidates for self-lubrication
Received 26 June 2013 at high temperatures, based on the formation of V2 O5 oxide. The aim of this investigation was to study
Received in revised form 18 October 2013 the effect of V additions on the structure, mechanical properties and oxidation resistance of Ti–Si–V–N
Accepted 19 October 2013
coatings deposited by DC reactive magnetron sputtering. The results achieved for TiSiVN films were com-
Available online 28 October 2013
pared and discussed in relation to TiN and TiSiN films prepared as reference. All coatings presented a fcc
NaCl-type structure. A shift of the diffraction peaks to higher angles with increasing Si and V contents
Keywords:
suggested the formation of a substitutional solid solution in TiN phase. Hardness and Young’s modulus
TiSi(V)N films
Structure
of the coatings were similar regardless on V content. The onset of oxidation of the films decreased sig-
Mechanical properties nificantly to 500 ◦ C when V was added into the films; this behaviour was independent of the Si and V
Oxidation resistance contents. The thermogravimetric isothermal curves of TiSiVN coatings oxidized at temperatures below
Vanadium oxide the melting point of ␣-V2 O5 (∼685 ◦ C) showed two stages: at an early stage, the weight increase over
time is linear, whilst, in the second stage, a parabolic evolution can be fitted to the experimental data. At
higher temperatures only a parabolic evolution was fitted. ␣-V2 O5 was the main phase detected at the
oxidized surface of the coatings. Reduction of ␣-V2 O5 to -V2 O5 phase occurred for temperatures above
its melting point.
© 2013 Elsevier B.V. All rights reserved.
0169-4332/$ – see front matter © 2013 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.apsusc.2013.10.117
F. Fernandes et al. / Applied Surface Science 289 (2014) 114–123 115
Table 1
Sample designation and deposition parameters of the coatings.
Sample Target power density (W/cm2) Deposition time (min) Pellets of V at the Ti target
Ti TiSi2
Interlayer Ti 10 – 5 –
Interlayer TiN – 15 –
[27–29]. Since the addition of V successfully decreased the friction substrates. Prior to the depositions, all the substrates were ultra-
coefficient of binary and ternary systems (down to reported values sonically cleaned in acetone for 15 min and alcohol for 10 min. The
of 0.2–0.3 at temperatures between 550 and 700 ◦ C), similar studies substrates were mounted in a substrate holder (which revolved
on the effect of V-addition should be performed on TiSiN coatings with 18 rev/min around the centre axis) giving a target to sub-
exhibiting unique mechanical properties and oxidation resistance. strate distance of 175 mm. Prior to deposition, the chamber was
This paper reports the first results on the effect of increasing evacuated down to 8.7 × 10−4 Pa and the substrates were etched
vanadium content to Ti–Si–V–N films deposited by DC reactive with Ar ion sputtering during 1 h with a bias voltage of −650 V to
magnetron sputtering. It is focused on coating structure, mechan- remove any surface contaminants. In order to enhance the adhesion
ical properties and oxidation resistance. Comparison of the results of the coatings, Ti and TiN adhesion layers of approximately 0.24
with those achieved for reference TiN and TiSiN coatings is pre- and 0.45 m, respectively, were deposited on the substrates before
sented. The results of this study will support further research aimed TiSi(V)N coatings. In all the depositions, the total working gas pres-
at the tribological behaviour of the coatings at high temperatures. sure was kept constant at 0.3 Pa, using approximately 30 sccm of
Ar and 17 sccm of N2 . The depositions were performed with a neg-
2. Experimental procedure ative substrate bias of 50 V. The deposition time was set in order to
obtain films with approximately 2.5 m of total thickness (includ-
Three series of TiSiN films with different Si contents, each ing interlayers). The deposition parameters are shown in Table 1.
series with increasing V contents (TiSiVN coatings), were deposited The chemical composition of the coatings was evaluated by
in a d.c. reactive magnetron sputtering machine equipped with electron probe microanalysis (EPMA—Cameca SX 50). Crystallo-
two rectangular (100 × 200 mm) magnetron cathodes working in graphic structure was investigated by X-ray diffraction (X’ Pert
unbalanced mode. A high purity Ti (99.9%) target, with 18 holes of Pro MPD diffractometer) using a grazing incidence angle of 1◦ and
10 mm in diameter (uniformly distributed throughout the prefer- Cu K␣1 radiation ( = 1.54060 Å). The XRD spectra were fitted by
ential erosion zone of the target), and a high purity TiSi2 (99.9%) using a pseudo-Voigt function to calculate either the full width
composite target were used in the depositions. The different sili- at half maximum (FWHM) and the peak position (2). The frac-
con contents were achieved by changing the power density applied ture cross-section morphology and the thickness of the films were
to each target. The V content was varied by changing the number investigated by scanning electron microscopy (SEM). Further, pro-
of high purity rods of Ti and V (with 10 mm in diameter) placed in filometer was used to confirm coating thickness and to evaluate
the holes of the Ti target. 4, 8 and 12 cylindrical pieces of vanadium surface roughness.
were used. In all the cases the total power applied to the targets was The hardness and Young’s modulus of films were measured in
set to 1500 W. To serve as reference, a stoichiometric TiN coating a nano-indentation equipment (Micro Materials NanoTest) using a
was deposited from the Ti target. Hereinafter the coatings will be Berkovich diamond pyramid indenter. In order to avoid the effect
designated as Sx-y, where x is related to the specific power applied of the substrate, the applied load (10 mN) was selected to keep
to the TiSi2 target (see Table 1), and therefore associated to the the indentation depth less than 10% of the coating’ thickness. 32
Si content on the coatings, and y the number of V rods used in measurements were performed in each sample. The level of resid-
the depositions, giving an indication of V content in each series of ual stresses was calculated through the Stoney equation [30] by
TiSiVN films. Thus, denomination S2-0, S2-4, S2-8 and S2-12 rep- measuring the bulk deflection of the film-substrate body.
resents coatings from the same series, i.e. with identical Si content The oxidation resistance of the coatings was evaluated by ther-
(TiSi2 target power 1.5 W/cm2 ), with increasing V content from 0 mogravimetric analysis (TGA) using industrial air (99.99% purity).
up to a maximum content achieved for the coating produced using The films deposited on alumina substrates were firstly heated with
12 vanadium pellets embedded into the Ti target. a constant temperature ramp of 20 ◦ C/min from room temperature
Polished high-speed steel (AISI M2) (Ø 20 × 3 mm, for mechani- up to 1200 ◦ C, in order to determine the onset point of oxidation.
cal properties measurements), FeCrAl alloy and alumina substrates Then, coated specimens of FeCrAl alloy were subjected to isother-
(10 × 10 × 1 mm, for oxidation tests and structural analysis), stain- mal tests at different selected temperatures and time. The weight
less steel discs (Ø 20 × 1 mm, for residual stress measurements) gain of the samples was evaluated at regular 2 s intervals using a
and (1 1 1) silicon samples (10 × 10 × 0.8 mm, for thickness mea- microbalance with an accuracy of 0.01 mg. The air flux used was
surements and chemical composition evaluation) were used as 50 ml/min and the heating rate up to the isothermal temperature
116 F. Fernandes et al. / Applied Surface Science 289 (2014) 114–123
Table 2
Chemical composition of the dissimilar coatings in at.%.
Coating at.%
N O Si Ti V
TiSiN S1-0 51.1 ± 0.6 0.4 ± 0.1 3.8 ± 0.0 44.7 ± 0.7 –
S1-4 51.4 ± 0.4 1.5 ± 0.3 3.1 ± 0.0 43.0 ± 0.7 1.1 ± 0.1
TiSiVN S1-8 51.5 ± 0.2 1.2 ± 0.1 2.9 ± 0.0 37.1 ± 0.3 7.3 ± 0.1
S1-12 49.7 ± 0.3 1.8 ± 0.1 2.8 ± 0.0 33.8 ± 0.3 12.0 ± 0.1
TiSiN S2-0 51.5 ± 0.2 0.6 ± 0.1 6.7 ± 0.0 41.3 ± 0.3 –
S2-4 52.1 ± 0.2 1.2 ± 0.1 5.7 ± 0.0 39.3 ± 0.4 1.6 ± 0.2
TiSiVN S2-8 51.7 ± 0.1 1.4 ± 0.1 5.6 ± 0.1 33.6 ± 0.3 7.6 ± 0.2
S2-12 51.1 ± 0.6 1.9 ± 0.2 5.3 ± 0.0 30.5 ± 0.6 11.3 ± 0.1
TiSiN S3-0 51.4 ± 0.4 1.8 ± 0.1 12.6 ± 0.1 34.2 ± 0.5 –
S3-4 52.9 ± 0.1 1.5 ± 0.1 10.7 ± 0.0 32.1 ± 0.2 2.7 ± 0.1
TiSiVN S3-8 52.8 ± 0.2 2.0 ± 0.1 10.7 ± 0.1 27.2 ± 0.2 7.3 ± 0.1
S3-12 52.4 ± 0.3 2.1 ± 0.2 10.3 ± 0.1 24.9 ± 0.9 10.4 ± 0.1
vanadium additions displayed lower hardness and Young’s modu- isothermal curves of TiSiN coating with V additions (S2-8), it can
lus values as compared to the other V rich coatings. The hardness be concluded that, independently of the isothermal temperature
enhance with V additions is probably due to the presence of V in and time of exposure, their mass gain is always much higher than
solid solution. In fact, the similar level of residual stresses as func- TiN and V-free TiSiN coatings (S2-0). The isothermal curves tested
tion of V content measured, which revealed to be from the range at 550 and 600 ◦ C showed two steps: at an early stage, the weight
of 3–4 GPa, and the observed shift of peaks to higher angles with V gain is rapidly increasing almost linearly up to approximately
incorporation, supports the previous affirmation. 0.1 mg/cm2 (particularly at 600 ◦ C), whereas in the second stage,
a parabolic evolution can be fitted. Isothermal annealing of S2-8
at 700 ◦ C shows significant increase of mass gain following the
3.2. Continuous and isothermal oxidation in air
parabolic evolution.
Fig. 3. (a) Effect of Si content on hardness and Young’s modulus for the TiN system.
(b) Effect of V content on hardness and Young’s modulus for coating S1-0.
Fig. 2. SEM fracture cross section micrographs of: (a) coating TiN, (b) coating S1-0,
(c) coating S1-12.
the white oxide was rich in Ti (crystals of TiO2 ), and the dark grey
phase composed mainly by Si and Ti. The Raman spectra of these
zones are plotted in Fig. 9. As can be observed the light grey phase
displays Raman active modes at 152, 253, 454 and 610 cm−1 [3,23]
assigned to rutile, confirming its presence in the oxidized coating
detected by XRD. The same phase, with much less intensity, was
detected for the dark grey phase; however, strong additional peaks
assigned to anatase (TiO2 ) were observed. This finding corroborates
the report of Pilloud et al. [23], who showed that anatase Raman
bands increased with the silicon content in TiSiN films, whereas
that of rutile decreased. As only strong peaks of rutile were detected
by XRD, this indicates that the amount of anatase should be small
and, therefore, not detectable by XRD diffraction. Since Raman pen-
etration depth is relatively limited, the absence of silicon oxide in
Raman spectra indicated that Si–O was below theTiO2 phase iden-
tified by XRD. This finding corroborates the results of Kacsich et al.
Fig. 5. Thermo gravimetric isothermal analysis of coatings exposed at different
temperatures.
[34,35]; they observed titanium oxide layer formed at outmost sur-
face and silicon oxide sublayer (TiSiN coatings).
The analysis of the oxidized surface of S2-8 coating reveals that
the surface oxide morphology is different from that of coating S2-
0. It displays a floret-like structure formed by light and dark grey
zones. At 550 ◦ C these phases are tiny distributed throughout the
surface, being difficult to clearly identify the boundaries between
them. At a temperature of 600 ◦ C, the separation between both
zones is evident with the shape of darker zone suggesting some type
of dendritic growth. With the increase of the isothermal time, an
increase in the area covered by the grey dark phase was observed.
The temperature 700 ◦ C led to a different oxide morphology, i.e.
the appearance of a black zone in some rosettes and the segrega-
tion of a white phase to the boundaries of the floret-like structure.
This change in the microstructure can be associated to the melting
of ␣-V2 O5 , which originates a smoother surface (Fig. 8d)). Fig. 10
shows examples of EDS spectra of points identified in Fig. 8 (phases
1 and 2). As can be observed, Ti K peak (4.931 keV) overlaps the V
K␣ peak (4.952 keV) and, therefore, V K peak should be taken in
Fig. 6. X-ray diffraction patterns of oxidized surface of: coating S2-8 annealed at
temperatures ranging from 550 to 700 ◦ C.
consideration for analyzing V importance. In order to identify the
different oxide phases marked in Fig. 8, the ratio between the peaks
intensities of V K (Si K␣) and Ti K␣ from EDS analyses are sum-
are shown in Figs. 7 and 8, respectively. EDS and Raman analyses marized in Table 3. To a relative increase in this ratio corresponds
were carried out on the oxidized surface of the coatings to identify the preferential formation of the oxide of that element. EDS analy-
their composition. On the oxidized surface of the S2-0 film, two dif- ses of points 1 and 2 reveal similar compositions for the grey light
ferent zones could be detected, a dark grey phase evenly distributed and dark zones, suggesting the presence of oxides containing Ti, Si
throughout the surface and white islands. EDS analysis showed that and V. However, darker phase has clearly a higher V content (see
Table 3). The Raman spectrum taken at grey dark phase (see Fig. 11)
shows the presence of intense peaks assigned to ␣-V2 O5 [3,8,10]
and small peaks related to rutile (TiO2 ). On light grey phase the
spectrum is much less defined, with broader bands. ␣-V2 O5 peaks
almost vanished whereas rutile bands are enhanced. Low crystal-
lized Ti–O phases give very similar Raman spectra, which evolve,
after annealing, either as anatase or rutile [41,42] phases. In sum-
mary, the dark grey phase can be assigned mainly to ␣-V2 O5 , and
small quantities of TiO2 , probably as a bilayer with ␣-V2 O5 on the
top and TiO2 underneath, as suggested by EDS measurements. In
fact, the high ratio of V/Ti of point 3 marked in Fig. 8, allows iden-
tifying a V-rich oxide, on the surrounding regions of the dark grey
phase. The light grey phase should be attributed to low crystallized
Ti(V)O2 . These results agree to the XRD results acquired on the oxi-
dized surface of the films where both Ti(V)O2 and ␣-V2 O5 phases
were indexed. When the isothermal time was prolonged from 10 to
30 min at 600 ◦ C, the increase of the intensity of the ␣-V2 O5 peak
observed in the XRD patterns was is in a good agreement with a
higher amount of grey dark phase observed on the oxidized surface
morphology.
Similarly to S2-0 coating signals from Si–O oxide were neither
Fig. 7. SEM observation of surface morphology of coating S2-0, after 1 h oxidation
detected by XRD nor by Raman spectroscopy. However, according
at 900 ◦ C. to EDS analysis this oxide should be present. As these signals are
120 F. Fernandes et al. / Applied Surface Science 289 (2014) 114–123
Fig. 8. Typical surface morphology of oxidized coatings: (a) coating S2-8 exposed to 550 ◦ C during 1 h, (b) coating S2-8 exposed to 600 ◦ C during 10 min, (c) coating S2-8
exposed to 600 ◦ C during 30 min, (d) coating S2-8 exposed to 700 ◦ C during 10 min.
Table 3
Ratio between the peaks intensities of V K (Si K␣) and Ti K␣ from EDS analyses of points marked in Fig. 8.
Points
1 2 3 4 5 6
4. Conclusion [12] A. Glaser, S. Surnev, F.P. Netzer, N. Fateh, G.A. Fontalvo, C. Mitterer, Oxida-
tion of vanadium nitride and titanium nitride coatings, Surf. Sci. 601 (2007)
1153–1159.
This investigation concerned the influence of V additions on [13] M. Uchida, N. Nihira, A. Mitsuo, K. Toyoda, K. Kubota, T. Aizawa,
the structure, mechanical properties and oxidation resistance of Friction and wear properties of CrAlN and CrVN films deposited by
Ti–Si–V–N coatings deposited by DC reactive magnetron sput- cathodic arc ion plating method, Surf. Coat. Technol. 177–178 (2004)
627–630.
tering. These coatings were compared to TiN and TiSiN films. [14] J.H. Ouyang, S. Sasaki, Tribo-oxidation of cathodic arc ion-plated
According to XRD analyses, all coatings showed an fcc NaCl-type (V,Ti)N coatings sliding against a steel ball under both unlubricated
structure assigned to crystalline TiN. A shift of the peaks to the and boundary-lubricated conditions, Surf. Coat. Technol. 187 (2004)
343–357.
right was observed with Si and V additions, indicative of a substitu-
[15] J.-K. Park, Y.-J. Baik, Increase of hardness and oxidation resistance of VN coat-
tional solid solution. Hardness and Young modulus of TiSiN coatings ing by nanoscale multilayered structurization with AlN, Mater. Lett. 62 (2008)
was insignificantly changed with increasing V content. The onset of 2528–2530.
[16] R. Franz, J. Neidhardt, R. Kaindl, B. Sartory, R. Tessadri, M. Lechthaler, P. Polcik,
oxidation of the coatings decreased with V additions down to tem-
C. Mitterer, Influence of phase transition on the tribological performance of arc-
peratures as low as 500 ◦ C, independently of the Si and V content in evaporated AlCrVN hard coatings, Surf. Coat. Technol. 203 (2009) 1101–1105.
the coatings. TiN and TiSiN coating exhibits a typical parabolic oxi- [17] Y. Qiu, S. Zhang, J.-W. Lee, B. Li, Y. Wang, D. Zhao, Self-lubricating CrAlN/VN
dation weight gain as a function of time, while a different evolution multilayer coatings at room temperature, Appl. Surf. Sci. 279 (2013) 189–196.
[18] W. Tillmann, S. Momeni, F. Hoffmann, A study of mechanical and tribological
is displayed by TiSiVN films. At temperatures below the melting properties of self-lubricating TiAlVN coatings at elevated temperatures, Tribol.
point of ␣-V2 O5 (∼685 ◦ C) two stages were exhibited: at an early Int. 66 (2013) 324–329.
stage, the weight increase over time is linear, whilst, in a second [19] Q. Luo, Temperature dependent friction and wear of magnetron sputtered coat-
ing TiAlN/VN, Wear 271 (2011) 2058–2066.
stage a parabolic evolution could be fitted to the experimental data; [20] S.H. Kim, J.K. Kim, K.H. Kim, Influence of deposition conditions on the
on the other hand, at high temperatures only a parabolic evolution microstructure and mechanical properties of Ti–Si–N films by DC reactive mag-
was fitted. ␣-V2 O5 showed to be the main phase present at the netron sputtering, Thin Solid Films 420–421 (2002) 360–365.
[21] M. Nose, Y. Deguchi, T. Mae, E. Honbo, T. Nagae, K. Nogi, Influence of
oxidized surface of coatings. Reduction of this phase occurred for sputtering conditions on the structure and properties of Ti–Si–N thin films
temperatures above their melting point. The relative amounts of prepared by r.f.-reactive sputtering, Surf. Coat. Technol. 174-175 (2003)
V2 O5 detected at the oxidized surface of V rich films are promising 261–265.
[22] S. Veprek, H.D. Männling, P. Karvankova, J. Prochazka, The issue of the
to achieve the envisaged good tribological properties; however it
reproducibility of deposition of superhard nanocomposites with hardness of
can be significantly compromised by their low oxidation resistance. ≥50 GPa, Surf. Coat. Technol. 200 (2006) 3876–3885.
[23] D. Pilloud, J.F. Pierson, M.C. Marco de Lucas, A. Cavaleiro, Study of the structural
changes induced by air oxidation in Ti–Si–N hard coatings, Surf. Coat. Technol.
Acknowledgments 202 (2008) 2413–2417.
[24] M. Diserens, J. Patscheider, F. Levy, Mechanical properties and oxidation resis-
This research is sponsored by FEDER funds through tance of nanocomposite TiN–SiNx physical-vapor-deposited thin films, Surf.
Coat. Technol. 120 (1999) 158–165.
the program COMPETE–Programa Operacional Factores de [25] F. Vaz, L. Rebouta, P. Goudeau, J. Pacaud, H. Garem, J.P. Rivière, A. Cavaleiro, E.
Competitividade–and by national funds through FCT – Fundação Alves, Characterisation of Ti1−x Six Ny nanocomposite films, Surf. Coat. Technol.
para a Ciência e a Tecnologia, under the projects: PEst- 133–134 (2000) 307–313.
[26] M. Diserens, J. Patscheider, F. Lévy, Improving the properties of titanium nitride
C/EME/UI0285/2013, CENTRO-07-0224-FEDER-002001 (Mais
by incorporation of silicon, Surf. Coat. Technol. 108–109 (1998) 241–246.
Centro SCT 2011 02 001 4637), PTDC/EME-TME/122116/2010 [27] Y.H. Cheng, T. Browne, B. Heckerman, E.I. Meletis, Mechanical and tribological
and Plungetec, as well as the grant (SFRH/BD/68740/2010). properties of nanocomposite TiSiN coatings, Surf. Coat. Technol. 204 (2010)
2123–2129.
[28] J. Patscheider, T. Zehnder, M. Diserens, Structure–performance relations in
References nanocomposite coatings, Surf. Coat. Technol. 146–147 (2001) 201–208.
[29] D. Ma, S. Ma, K. Xu, The tribological and structural characterization of nano-
[1] Z. Zhou, W.M. Rainforth, D.B. Lewis, S. Creasy, J.J. Forsyth, F. Clegg, A.P. Ehiasar- structured Ti–Si–N films coated by pulsed-d.c. plasma enhanced CVD, Vacuum
ian, P.E. Hovespian, W.D. Münz, Oxidation behaviour of nanoscale TiAlN/VN 79 (2005) 7–13.
multilayer coatings, Surf. Coat. Technol. 177–178 (2004) 198–203. [30] G. Stoney, The tension of metallic films deposited by electrolysis, Proc. R. Soc.
[2] M. Pfeiler, K. Kutschej, M. Penoy, C. Michotte, C. Mitterer, M. Kathrein, The effect London, Ser. A 82 (1909) 172–175.
of increasing V content on structure, mechanical and tribological properties of [31] C.P. Constable, D.B. Lewis, J. Yarwood, W.D. Münz, Raman microscopic studies
arc evaporated Ti–Al–V–N coatings, Int. J. Refract. Met. Hard Mater 27 (2009) of residual and applied stress in PVD hard ceramic coatings and correlation
502–506. with X-ray diffraction (XRD) measurements, Surf. Coat. Technol. 184 (2004)
[3] N. Fateh, G.A. Fontalvo, G. Gassner, C. Mitterer, Influence of high-temperature 291–297.
oxide formation on the tribological behaviour of TiN and VN coatings, Wear [32] F. Vaz, L. Rebouta, B. Almeida, P. Goudeau, J. Pacaud, J.P. Rivière, J. Bessa e Sousa,
262 (2007) 1152–1158. Structural analysis of Ti1−x Six Ny nanocomposite films prepared by reactive
[4] E. Badisch, G.A. Fontalvo, M. Stoiber, C. Mitterer, Tribological behavior of PACVD magnetron sputtering, Surf. Coat. Technol. 120–121 (1999) 166–172.
TiN coatings in the temperature range up to 500 ◦ C, Surf. Coat. Technol. 163–164 [33] M. Diserens, J. Patscheider, F. Lévy, Mechanical properties and oxidation resis-
(2003) 585–590. tance of nanocomposite TiN–SiNx physical-vapor-deposited thin films, Surf.
[5] M. Stoiber, E. Badisch, C. Lugmair, C. Mitterer, Low-friction TiN coatings Coat. Technol. 120–121 (1999) 158–165.
deposited by PACVD, Surf. Coat. Technol. 163–164 (2003) 451–456. [34] T. Kacsich, S. Gasser, Y. Tsuji, A. Dommann, M.A. Nicolet, A. Nicolet, Wet oxida-
[6] E. Lugscheider, O. Knotek, K. Bobzin, S. Bärwulf, Tribological properties, phase tion of Ti34 Si23 B43 , Br. J. Appl. Phys. 85 (1999) 1871–1875.
generation and high temperature phase stability of tungsten- and vanadium- [35] T. Kacsich, M.A. Nicolet, Moving species in Ti34 Si23 N43 oxidation, Thin Solid
oxides deposited by reactive MSIP–PVD process for innovative lubrication Films 349 (1999) 1–3.
applications, Surf. Coat. Technol. 133–134 (2000) 362–368. [36] Z. Zhou, W.M. Rainforth, C. Rodenburg, N.C. Hyatt, D.B. Lewis, P.E. Hovsepian,
[7] D.B. Lewis, S. Creasey, Z. Zhou, J.J. Forsyth, A.P. Ehiasarian, P.E. Hovsepian, Q. Oxidation Behavior and mechanisms of TiAlN/VN coatings, Metall. Mater. Trans.
Luo, W.M. Rainforth, W.D. Münz, The effect of (Ti + Al):V ratio on the structure A 38 (2007) 2464–2478.
and oxidation behaviour of TiAlN/VN nano-scale multilayer coatings, Surf. Coat. [37] R. Franz, J. Neidhardt, C. Mitterer, B. Schaffer, H. Hutter, R. Kaindl, B. Sartory,
Technol. 177–178 (2004) 252–259. R. Tessadri, M. Lechthaler, P. Polcik, Oxidation and diffusion processes during
[8] K. Kutschej, P.H. Mayrhofer, M. Kathrein, P. Polcik, C. Mitterer, Influence of annealing of AlCrVN hard coatings, J. Vac. Sci. Technol., A 26 (2008) 302–308.
oxide phase formation on the tribological behaviour of Ti–Al–V–N coatings, [38] M. Wittmer, J. Noser, H. Melchior, Oxidation kinetics of TiN thin films, J. Appl.
Surf. Coat. Technol. 200 (2005) 1731–1737. Phys. 52 (1981) 6659–6664.
[9] K. Kutschej, P.H. Mayrhofer, M. Kathrein, P. Polcik, C. Mitterer, A new low- [39] C.W. Zou, X.D. Yan, J. Han, R.Q. Chen, W. Gao, Microstructures and optical prop-
friction concept for Ti1−x Alx N based coatings in high-temperature applications, erties of -V 2 O 5 nanorods prepared by magnetron sputtering, J. Phys. D: Appl.
Surf. Coat. Technol. 188–189 (2004) 358–363. Phys. 42 (2009) 145402.
[10] G. Gassner, P.H. Mayrhofer, K. Kutschej, C. Mitterer, M. Kathrein, A new low [40] A. Bouzidi, N. Benramdane, S. Bresson, C. Mathieu, R. Desfeux, M.E. Marssi, X-ray
friction concept for high temperatures: lubricious oxide formation on sputtered and Raman study of spray pyrolysed vanadium oxide thin films, Vib. Spectrosc.
VN coatings, Tribol. Lett. 17 (2004) 751–756. 57 (2011) 182–186.
[11] P.H. Mayrhofer, P.E. Hovsepian, C. Mitterer, W.D. Münz, Calorimetric evidence [41] M. Fernández-García, X. Wang, C. Belver, J.C. Hanson, J.A. Rodriguez, Anatase-
for frictional self-adaptation of TiAlN/VN superlattice coatings, Surf. Coat. Tech- TiO2 nanomaterials: morphological/size dependence of the crystallization and
nol. 177–178 (2004) 341–347. phase behavior phenomena, J. Phys. Chem. C 111 (2006) 674–682.
F. Fernandes et al. / Applied Surface Science 289 (2014) 114–123 123
[42] R.J. Gonzales, Raman, Infrared, X-ray, and EELS Studies of Nanophase Titania, [44] P. Balog, D. Orosel, Z. Cancarevic, C. Schon, M. Jansen, V2 O5 phase diagram
in: Ph.D. Dissertation, Virginia Polytechnic Institute, Blacksburg, Virginia, 1996, revisited at high pressures and high temperatures, J. Alloys Compd. 429 (2007)
pp. 426. 87–98.
[43] V.P. Filonenko, M. Sundberg, P.E. Werner, I.P. Zibrov, Structure of a high- [45] J.G. Keller, D.L. Douglass, The high-temperature oxidation behavior of
pressure phase of vanadium pentoxide beta-V2 O5 , Acta Crystallogr., Sect. B: vanadium–aluminum alloys, Oxid. Met. 36 (1991) 439–464.
Struct. Sci. 60 (2004) 375–381.
Annex H
Annex H
F. Fernandes, J. Morgiel, T. Polcar, A. Cavaleiro, Oxidation and diffusion
processes during annealing of TiSi(V)N films, (2014), under review, “Thin Solid
Films”.
171
Filipe Daniel Fernandes
Annex H
1
SEG-CEMUC - Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis
Santos, 3030-788 Coimbra, Portugal.
2
Institute of Metallurgy and Materials Science of Polish Academy of Sciences, Krakow, Poland
3
National Centre for Advanced Tribology (nCATS), School of Engineering Sciences, University of
Southampton, Highfield, Southampton, SO17 1BJ, UK.
4
Department of Control Engineering Czech Technical University in Prague Technicka 2, Prague 6, 166
27 Czech Republic.
*Email address: filipe.fernandes@dem.uc.pt, tel. + (351) 239 790 745, fax. + (351) 239 790 701
Abstract
The effect of V additions on oxidation resistance, oxide scale formation and diffusion
processes for TiSiVN system and their comparison to TiSiN is investigated. A dual layer
oxide was formed in the case of TiSiN coating with a protective Si-O layer at an
oxide/coating interface; however, in zones of film defects a complex oxide structure was
developed. V additions increased the oxidation rate of the coatings as a result of the inversion
of the oxidation mechanism due to ions diffusion throughout the oxide scale, which inhibited
the formation of a continuous protective silicon oxide layer.
Keywords: TiSiVN system, structural evolution, oxidation, oxide scale, diffusion processes
1. Introduction
173
Filipe Daniel Fernandes
Annex H
have attracted the scientific community and particular attention has been given to vanadium
oxide. The beneficial influence of Magnéli oxides formed by oxidation of vanadium on the
friction has already been reported for ternary CrVN [2], (V,Ti)N [3], multilayer AlN/VN [4],
quaternary single layer or multilayered CrAlVN [5-6], and TiAlVN [7-8] systems. Recently
we have reported the effect of V incorporation on the structure, mechanical properties and
oxidation resistance of TiSiVN films deposited by DC reactive magnetron sputtering [9].
Lubricious vanadium oxides have been successfully detected on the oxidized surface of these
films. Here we show the thermal stability, oxide formation and, particularly, diffusion
processes during annealing of TiSiVN films deposited by magnetron sputtering.
2. Experimental procedure
TiSiN and TiSiVN coatings with approximately the same silicon content and about 2.5
µm of total thickness, were deposited on FeCrAl alloy substrates by a d.c. reactive magnetron
sputtering machine equipped with two rectangular, Ti (99.9%) and TiSi2 (99.9%), magnetron
cathodes working in unbalanced mode. V incorporation was achieved by inserting 8 pellets of
vanadium into the erosion zone of Ti target. In both cases Ti and TiN adhesion layers were
deposited as bonding layers improving coating to substrate adhesion. The interlayer also
contained V when V pellets are placed in the Ti target. A detail description of the deposition
conditions is reported elsewhere [9]. Temperature effect on the structure of the V rich coating
was characterized by in-situ hot-XRD analysis in the range from 500 ºC up to 750 ºC, in open
air, using a grazing incidence angle of 2º and Co Kα radiation (1.789010 Å). Between each
selected temperature a step of 10 min holding time was allowed for thermal stabilization and
30 min time acquisition was used. Oxidation of films was assessed by thermogravimetric
analysis (TGA) using industrial air (99.99% purity). The coatings were isothermal tested at
different selected temperatures and times, based on the thermogravimetric oxidation curves of
films performed at constant linear-temperature ramp (RT to 1200◦C at 20◦C/min) shown in
reference [9]. These isothermal tests were performed having in consideration the temperatures
where the oxidation showed takes mainly place. After annealing, the cross section thin foils of
oxidized films was prepared by focused ion beam (FIB) and analyzed by transmission
electron microscope (TEM) equipped with an energy-dispersive x-ray (EDS) spectroscopy
system.
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RT
Relative Intensity
(after
cooling down)
0
750 C
0
700 C
0
675 C
0
650 C
0
600 C
0
550 C
0
500 C
RT
2
16 18 20 22 24 26 28 30 32 34 36 38 40 42 44 46 48 50 52 54
TiN
Ti(V)O2
V
V
'V
V3O5
V2 O 3
VO
Figure 1 - XRD spectra of TiSiVN film at different temperatures (upper) and position of main peaks of
corresponding phases (lower).
TiSiN film was annealed at 900 0C for 1 hour, TiSiVN coating was annealled at two
different isothermal temperatures. The results of the thermo gravimetric analysis are presented
in Fig. 2 a). As expected, TiSiN film is much more resistant to oxidation, showing a typical
parabolic oxidation weight gain as a fuction of time, which indicates the presence of
protective oxide scales. Vanadium addition to TiSiN film strongly reduced the oxidation
resistance. The isothermal curves at 550 ºC and 600 ºC showed two steps: they started with a
linear increase in mass gain and then followed with a parabolic evolution. The surface
morphologies of the oxidized coatings are shown in Fig. 2 b-c) for coatings TiSiN tested at
900 ºC and TiSiVN oxidized at 600 ºC, respectively. The detailed description of the surface
oxide constitution, based on XRD diffraction, Raman spectroscopy and SEM-EDS analyses,
can be found in our previous study [9]. Therefore we will only summarize here the main
results to support investigation aimed at diffusion processes and described later. Annealed
TiSiN film (see Fig. 2b) displayed two different surface features: white and dark gray islands
evenly distributed on the surface. Raman analyses revealed that white phase was rutile (TiO 2),
while dark gray phase was a mixture of rutile and anatase. However, only rutile peaks were
detected by XRD diffraction suggesting very limited amount of anatase in the dark gray
islands. Furthermore, strong signals of Si were detected on dark zone by EDS. However, the
signals of silicon oxide were neither detected by XRD nor by Raman spectroscopy indicating
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amorphous character, which corroborates previous reports [13-14]. Silicon oxide was
positioned below Ti-O rich layer. Different morphology and oxide phases were detected on
the oxidized surface of TiSiVN coating. At 600 ºC, the film displayed a floret-like structure
formed by light and dark gray phases. XRD diffraction showed the presence of α-V2O5 and
Ti(V)O2 oxides. Although EDS analysis revealed similar spectra for both phases, the much
less intensity of the V peak on the light gray phase, and its conjugation with Raman analysis
allowed identifying dark and light gray phases as α-V2O5 and Ti(V)O2 oxides, respectively.
Similar phases were detected at 550 ºC by XRD; however, only small dark gray areas were
observed at the surface. These results match well with high temperature in-situ XRD
diffractograms (Fig. 1), where these phases were indexed. Similar to TiSiN coating, the signal
from Si-O phase was neither detected by XRD nor by Raman.
0.20
S2-8 600 ºC 30 min
Weight gain (mg/cm )
2
0.15
S2-8 500 ºC 1 h
0.10
S2-0 900 ºC 1 h
0.05
0.00
0 1200 2400 3600
a) Time (s)
Gray phase
5 µm 20 µm
Dark gray phase
White phase Dark phase
50 µm 500 µm
b) c)
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Figure 3 shows the growing mechanism of the V2O5 phase over the Ti(V)O2 oxide.
Nucleation points of V2O5 started being formed over the oxidized surface, growing up by
vanadium lateral diffusion as suggested by the V depleted white and V rich black regions
marked in the SEM micrograph. In order to investigate the oxide scale films growth, TEM
cross-sections were prepared by FIB from TiSiN and TiSiVN films annealed at 900 and 600
ºC, respectively.
Floret-like V2O5 phase
V rich black
region
Figure 3 - Oxidized surface of TiSiVN coating showing the growing mechanism of V2O5 phase.
Fig. 4 and 5 displays the bright-field STEM images, associated elemental maps and
scanning elemental profiles from cross section of oxidized TiSiN and TiSiVN coatings. In the
case of TiSiN film a multilayer of oxides can be identified, being more complex close to the
film defect (white zones) shown in Fig. 4a): (i) an outer Ti-rich layer comprised by shaped
crystals with bigger size in the top of the film defect, zone corresponding to the white phase
and, (ii) a Si-rich layer, which is itself divided in 3 layers on the zone around the defect, an
intermediate layer containing Ti, sandwiched between two Ti-free layers, being the external
porous and the internal one very compact. Far from the film defect (left zone of the
micrograph), below to the TiO2 crystals only a homogeneous Si-rich layer was observed. The
measured elemental cross-section depth profiles for two lines were plotted in Fig. 4b) and 4c),
respectively. The profiles corroborated STEM/EDX elemental mapping showing in detail the
composition of surface oxides. The surface layer formed exclusively of Ti-O was followed by
a Si-rich layer with scattering in the signals intensity, in agreement with brightness intensity
in figures 4 a). It is clear in Ti-signal a small increase in intensity in the zone of high-Si
content. This variation is more intense in line 2 than in line 1, suggesting an influence of the
film defect, which is closer to line 2. Continuous and compact Si-O layer at the oxide/coating
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interface acts as an efficient barrier against oxygen and metal ions diffusion and thus
protecting the coating from further oxidation [13-14].
It is evident that Ti oxide forms first on the surface of the film due to the higher
affinity of Ti for O than Si [13-14] and then, with further increase in temperature, silicon
oxide is formed due to the progressive segregation of Si. The oxidation process will then
occur through the inwards diffusion of O2- through the TiO2 layer and the outwards diffusion
of Ti4+ through the Si-O layer [15]. Nevertheless, due to the presence of film defects, in the
first stage of the oxidation process, a high amount of Ti ions will be supplied in that zone
corresponding to the oxidation occurring in the defect walls. This process will promote the
formation of a porous Si-O layer, with low diffusion barrier performance to the out diffusion
of Ti4+ ions, leading to large TiO2 crystals on the oxide scale surface. From the moment that a
Si-O barrier layer is formed in the defect walls (see Si signal in figure 4 a)) the oxidation
process will be controlled by the O2- inward and Ti4+ outward ions diffusion trough the Si-O
layer [15]. It should be remarked that this phenomenon should not influence significantly the
global oxidation behavior of the Ti-Si-N coating (which shows the parabolic behavior
presented in Fig. 2 a), in agreement with literature [13-15]) since, on the one hand, it only
occurs in a few defects in the films surface (see Fig. 2 b)) and, on the other, it should be more
intense in the first stage of the oxidation, during heating up to the 900 ºC isothermal oxidation
temperature.
Gray phase
White phase
Line 2
Line 1
STEM O Si Ti
2 µm
a)
Figure 4 - a) Bright field STEM/EDX maps of TiSiN coating oxidized at 900 0C during 1h. b) and c). Elemental
profiles along the cross section of the oxidized coating from: b) line 1, c) line 2.
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Ti K
400 O K / Ti K
Si K / Ti K
300
Substrate
200
100
Counts
ne
TiN interlayer
Ti interlayer
zo
d
0.8 ize Coating
xid
O
0.4
0.0
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5
b) Distance (m)
Ti K
200 O K / Ti K
ne Si K / Ti K
zo
150 d
ize
xid
Substrate
100 O
50
Counts
TiN interlayer
Ti interlayer
Coating
0.8
0.4
0.0
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5
c) Distance (m)
Figure 4 (continued).
TiSiVN coating exhibited, after annealing, two-layers oxide structure with a thick
porous inner layer and an outer discontinuous layer of well-defined crystals; the latter
corresponded to the gray areas described above. The elemental maps shown in Fig. 5a)
suggested that the majority of vanadium was located in the outer crystal layer, whereas inner
layer was Ti and Si rich. This result corroborates the detection of the rutile type compound
Ti(V)O2 indexed by XRD and Raman spectroscopy in the light gray phase in Fig. 2 [9].
Elemental lines shown in Fig. 5b) and 5c) showed that diffusion of V occurred exclusively
within the oxidized volume, since a constant signal was measured for the remaining non-
oxidized coating. In general, we observed a thicker non oxidized layer in the regions covered
by V-O crystal phase. Furthermore, comparing either the V content integrated intensities of
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Fig. 5a) and b) or the brightness of V signal between the right and left part of Fig. 5 a), it can
be concluded that V has to be diffused from the left to the right zone of this figure. Finally, it
should be remarked that there is no indication of formation of any Si-rich layer, being Si
signal uniformly distributed in both oxide scale and non-oxidized coating. Therefore, a dense
compact protective silicon oxide layer localized in subsurface was not formed being Si-O
randomly distributed in the Ti(V)O2 porous scale.
STEM Ti Si V
2 µm
a)
Ti K
2000 O K / Ti K
Si K / Ti K
1600
V K / Ti K
1200
Substrate
800
TiN interlayer
Counts
400
Ti interlayer
0.4
0.0
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5
b) Distance (m)
Figure 5 - a) Bright field STEM/EDX maps of TiSiVN coating oxidized at 600 0C during 30 min. Elemental
profiles along the cross section of the oxidized coating from: b) line 1, c) line 2.
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Ti K
2000 O K / Ti K
Si K / Ti K
1600
V K / Ti K
1200
Substrate
800
Increasing of V
TiN interlayer
Counts
400
Ti interlayer
and O counts
0.4
0.0
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5
c) Distance (m)
Figure 5 (continued).
Isothermal curve of TiSiVN coating oxidized at temperatures below the melting point
of V2O5 started with a linear increase in mass gain. This step, where a continuous Ti(V)O2
oxide layer was grown (see Fig. 1), should be attributed to the conjugation of a fast rate of
mass transport in the oxide with a slow surface step accompanying the gas dissociation,
adsorption and subsequent entry of oxygen into the Ti(V)O2 lattice. In fact, at the very
beginning of the oxidation process, due to the high Ti content, a TiO2 layer starts growing.
The presence of V ions and its high solubility with Ti promote the formation of Ti(V)O2 solid
solution, which comprises vanadium cations with lower oxidation states as V3+ ions [16-17].
The substitution of Ti4+ by V3+ ions in the TiO2 lattice would increase the concentration of
interstitial metallic Ti4+ ions and decrease the number of excess electrons. The oxidation
mechanism is inverted in relation to Ti-N, being the outwards diffusion of metallic ions (Ti3+
Tin+) more favorable that the inwards diffusion of O2- ions. There are several consequences of
this inversion: (i) the oxidation rate increases substantially, (ii) the kinetics of the oxidation is
not controlled any more by a diffusion process but by the rate of dissociation, adsorption and
combination at the surface of O2- with metallic ions, and (iii) continuous and compact Si-rich
layer cannot be formed. In fact, the growing of Ti(V)-O on the external part of the scale
leaves a less compact zone in the interface between the non-oxidized material and the oxide
scale, facilitating the local segregation of Si inside the Ti(V)-O instead its segregation for the
interface. The elemental maps suggest that Si-O and Ti(V)O2 coexist in the inner porous
layer, being unable to protect the material from oxidation. As a consequence, the oxidation
process is accelerated, which is demonstrated by the initial rapid oxidation observed in the
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isothermal curve. Thongtem et al. [16] studied the effect of V doping to Ti-based alloys and
observed similar behavior with increasing V content.
The experimental evidence shows that V2O5 crystals were formed exclusively on the
surface. When V3+ or V4+ arrive to the surface, they are further oxidized to V5+ combining
with O ions and creating nucleation points for the V2O5 growth (see Fig. 3), which will
expand and grow up by vanadium lateral diffusion. The process develops leading to the
formation of the observed floret-like dispersive V2O5 phases that can comprise both α-V2O5
and β-V2O5 phases as evidenced in the XRD patterns at high temperature. As soon as V2O5
oxide expands laterally, the surface will be progressively covered by this oxide making more
and more difficult the access of both types of ions, up to the moment that the rate of
dissociation, adsorption and combination of O2- with metallic ions is not anymore the kinetics
controlling step but the ion diffusion through the oxide scale. Therefore, after this moment the
isothermal curve will follow a parabolic law (see Fig. 2a)). It should be remarked that V2O5 is
more protective than (Ti,V)O2 since the non-oxidized coating underneath it is thicker.
4. Conclusion
Acknowledgments
This research is sponsored by FEDER funds through the program COMPETE – Programa
Operacional Factores de Competitividade – and by national funds through FCT – Fundação
para a Ciência e a Tecnologia –, under the project: PTDC/EME-TME/122116/2010, as well as
the grant (SFRH/BD/68740/2010). T. Polcar acknowledges support from the project MSM
6840770038.
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References
[2] M. Uchida, N. Nihira, A. Mitsuo, K. Toyoda, K. Kubota, T. Aizawa, Friction and wear
properties of CrAlN and CrVN films deposited by cathodic arc ion plating method, Surf.
Coat. Technol., 177–178 (2004) 627-630.
[3] J.H. Ouyang, S. Sasaki, Tribo-oxidation of cathodic arc ion-plated (V,Ti)N coatings
sliding against a steel ball under both unlubricated and boundary-lubricated conditions, Surf.
Coat. Technol., 187 (2004) 343-357.
[4] J.-K. Park, Y.-J. Baik, Increase of hardness and oxidation resistance of VN coating by
nanoscale multilayered structurization with AlN, Mater. Lett., 62 (2008) 2528-2530.
[6] Y. Qiu, S. Zhang, J.-W. Lee, B. Li, Y. Wang, D. Zhao, Self-lubricating CrAlN/VN
multilayer coatings at room temperature, Applied Surface Science, 279 (2013) 189-196.
[7] Q. Luo, Temperature dependent friction and wear of magnetron sputtered coating
TiAlN/VN, Wear, 271 (2011) 2058-2066.
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[11] R. Franz, C. Mitterer, Vanadium containing self-adaptive low-friction hard coatings for
high-temperature applications: A review, Surf. Coat. Technol., 228 (2013) 1-13.
[13] T. Kacsich, S. Gasser, Y. Tsuji, A. Dommann, M.A. Nicolet, A. Nicolet, Wet oxidation
of Ti34Si23B43, Journal of Applied Physics, 85 (1999) 1871-1875.
[14] T. Kacsich, M.A. Nicolet, Moving species in Ti34Si23N43 oxidation, Thin Solid Films,
349 (1999) 1-3.
[17] A. Glaser, S. Surnev, F.P. Netzer, N. Fateh, G.A. Fontalvo, C. Mitterer, Oxidation of
vanadium nitride and titanium nitride coatings, Surface Science, 601 (2007) 1153-1159
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Annex I
F. Fernandes, T. Polcar, A. Cavaleiro, Tribological properties of self-lubricating
TiSiVN coatings at room temperature, (2014), accept for publication, “Surface and
Coatings Technology”, DOI: 10.1016/j.surfcoat.2014.10.016
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1
SEG-CEMUC - Department of Mechanical Engineering, University of Coimbra, Rua Luís
Reis Santos, 3030-788 Coimbra, Portugal.
2
Department of Control Engineering Czech Technical University in Prague Technicka 2,
Prague 6, 166 27 Czech Republic.
3
n-CATS University of Southampton Highfield Campus SO17 1BJ Southampton, UK.
*
Email address: filipe.fernandes@dem.uc.pt, tel. + (351) 239 790 745, fax. + (351) 239 790
701
Abstract
In the last years, vanadium rich coatings have been introduced as possible candidates
for self-lubrication due to their optimum tribological properties. In the present investigation,
the influence of V incorporation on the wear performance of TiSiN films deposited onto WC
substrates by d.c. reactive magnetron sputtering is reported. The results achieved for TiSiVN
films were compared and discussed in relation to Ti0.80Si0.15N, TiN and Ti0.82V0.15N coatings
prepared as references. The tribological properties of the coatings were evaluated at room
temperature on a pin on disc tribometer equipment using two different counterparts: Al2O3
and HSS balls. The wear tracks, ball-wear scars and wear debris were characterized by
scanning electron microscopy with energy dispersive X-ray spectroscopy (SEM-EDS).
Tribological tests indicated that the wear rate and the friction coefficient of Ti0.80Si0.15N
coating decreased with continuous increase of V content being the overall behaviour strongly
dependent on the counterpart ball material. For Al2O3 balls the wear rate and friction
coefficient of coatings were much lower compared to sliding against HSS steel balls.
Ti0.80Si0.15N showed the lowest wear resistance among all tested coatings, independently of
the counter-body. For V rich coatings tested with Al2O3 balls the polishing wear mechanism
was observed, whereas adhesion wear took place when tested against HSS balls.
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1. Introduction
Over the past years, ternary titanium-based nitride coatings, such as TiXN (X = B, Cr,
Al, Si, Cr, C, etc.), have been largely employed in industrial applications involving cutting
tools, wear protection, machinery components and high temperature parts, due to their high
hardness and suitable wear and oxidation resistance [1, 2]. However, in dry machining
operations, the tools experience extremely harsh conditions combining abrasion, adhesion,
oxidation, mechanical and thermal loads, which limit their lifetime. Since low friction
coefficients can effectively reduce the contact temperatures during sliding and, consequently,
improve the tribological behavior of tools, a lot of efforts has been carried out in the last years
to develop self-lubricant coatings. These films should retain beneficial properties of (TiX)N
films and offer lubricity through the formation of low-friction surface oxides. Recently, the
attention of the scientific community has been focused on the formation of thin reaction films
of the so-called Magnéli oxide phases based on Ti, Si, Mo, W and V, which possess easy
shear planes [3, 4]. Among these elements, particular attention has been given to vanadium-
containing coatings which form VnO3n-1 Magnéli phases. Over the whole ternary systems the
effect of V doping was only extensively studied for ternary CrAlN [5, 6] and TiAlN [2, 7-9]
coatings in single layer or multilayered configurations. In all the cases a beneficial influence
of V additions on the tribological behaviour of the coatings at room and high temperatures has
been demonstrated. Since other ternary coating systems with superior mechanical properties
have been successfully used as protective coatings in tribological applications, the effect of
vanadium doping on their structure should also be considered. This is the case of TiSiN
system, which shows similar levels of oxidation resistance as CrAlN and TiAlN films.
Preliminary studies on the effect of V additions on the structure, mechanical properties and
oxidation resistance of TiSiN systems have been carried out in our previous investigation
[10]. Hence, as a way to complement that study, the aim of this work was to characterize the
tribological behaviour at room temperature of TiSiVN sputtered coatings, sliding against
Al2O3 and HSS balls. The influence of V addition on the TiSiN system was compared with
the tribological results of TiN, Ti0.82V0.15N and Ti0.80Si0.15N coatings prepared and tested as a
reference.
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2. Experimental Procedure
TiSiVN coatings with dissimilar Si and V contents, evaluated in our previous work
[10], were deposited by d.c. reactive magnetron sputtering from two rectangular (100×200
mm) magnetron cathodes working in unbalanced mode. High speed steel (AISI M2), FeCrAl
alloy, cemented carbides, silicon and stainless steel substrates were coated for mechanical,
structural, chemical composition and residual stress measurements. A high purity Ti (99.9%)
target, with 10 mm diameter holes, and a high purity TiSi2 (99.9%) target were used in the
depositions. The coatings comprise three TiSiVN coatings with different V and Si contents
and three other coatings deposited as a reference: TiN, Ti0.82V0.15N and Ti0.80Si0.15N. V
incorporation was achieved by filling Ti holes with 4 and 8 V pellets being the remaining
empty holes filled with pellets of Ti. Silicon content was varied by changing the power
density applied to each target. Ti and TiN adhesion layers of approximately 0.24 and 0.45 µm,
respectively, were deposited below final TiSi(V)N coatings. The interlayers also contained V
when V pellets were placed in the Ti target. A negative bias of 50 V was applied to the
substrate holder. All the depositions were performed at a total working gas pressure of 0.3 Pa.
The deposition time was set in order to produce coatings with approximately 2.5 µm of total
thickness, including interlayer referred to above. A full description of the deposition
procedure is given elsewhere [10]. Summary of the deposition parameters used for the
coatings production and their chemical composition determined by electron probe
microanalysis (EPMA) are shown in Table 1 and Table 2, respectively. The denomination of
the samples, as presented in tables 1 and 2, will be adopted throughout the paper in order to
help the coatings identification and the paper reading. The subscript in Ti, Si and V letters
represents the ratio between the at.% of each element in the coating in relation to the at.% of
Nitrogen.
The structure of the films was analyzed by X-ray diffraction (X’ Pert Pro MPD
diffractometer) using a grazing incidence angle of 1º with Cu Kα1 radiation (λ = 1.54060 Å).
The adhesion of the films on the M2 steel substrate was evaluated by a scratch-test apparatus.
Specimens were scratched as the normal force was increased linearly from 5 to 70 N, using a
Rockwell C indenter with a spherical tip with a radius of 0.2 mm, a scratch speed of 10
mm/min and a loading speed of 100 N/min. Critical loads were determined by optical
microscope analysis according to the standard for scratch-test evaluation [11]. For each
specimen, the indicated critical loads results from the average of 3 different scratches. The
hardness and Young’s modulus of films were evaluated by depth-sensing indentation
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technique (Micro Materials NanoTest) using a Berkovich diamond pyramid indenter. In order
to avoid the effect of the substrate, the applied load (10 mN) was selected to keep the
indentation depth less than 10% of the coating´s thickness. The residual stress values were
calculated through the Stoney equation [12] by measuring the differential deflection between
an uncoated and coated stainless steel 25 mm diameter disk.
The tribological behaviour of the coatings was evaluated by dry-sliding tests at room
temperature on a pin on disc tribometer. Al2O3 and HSS steel balls (both with a diameter of 6
mm) were used as counterparts. The radius of the wear tracks was set to 5.3 and 4 mm for
Al2O3 and HSS steel balls, respectively. All the measurements were performed with a linear
speed of 0.10 m.s-1, load of 5 N, relative humidity 48±5% and 5000 cycles. The wear rate of
the coatings was determined from the area of the wear track cross section using a 3D optical
profilometer; an average from four different locations was used. The ball wear rates were
evaluated from the wear cap images taken in an optical microscope. To ensure the
reproducibility of results a set of three tests was performed under identical conditions for each
coating. After the wear tests, the wear tracks and wear debris were characterized by scanning
electron microscopy with energy dispersive X-ray spectrometry (SEM-EDS).
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The X-ray diffraction patterns of the as-deposited films are shown in Figure 1. All the
coatings presents an fcc NaCl-type structure assigned to crystalline TiN phase, but different
peaks position and intensities can be perceived. All films presented well defined peaks except
for the coating with higher silicon content (Ti0.52Si0.20V0.14N) for which a strong decrease of
the peaks intensity as well a significant broadening was observed. This result suggests a
decrease of the crystallinity and grain size, in good accordance with literature where for TiSiN
films with high Si contents amorphization of the structure is usually observed [13]. A closer
analysis of the peaks positions indicated that in all cases a fcc solid solution should be
expected. In fact, starting with the individual addition of V and Si to TiN, the peaks are
shifted to higher angles, behaviour associated to a smaller unit cell (insert in Figure 1). Taking
into account that the residual stresses are very similar among these coatings (see table 3), such
a result can only be explained by the presence of V or Si in substitutional solid solution, due
to their smaller radius (0.143 and 0.117 nm, respectively) in comparison with Ti (0.147 nm).
When both Si and V are added to TiN, a synergetic effect of both elements are observed being
the diffraction peaks shifted to further higher diffraction angles. A detailed description of the
effect of Si and V additions on the structure of TiN is shown in our previous work [10].
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TiN (111)
TiN (111)
Relative Intensity
Ti0.52Si0.20V0.14N
Ti0.65Si0.11V0.15N
Ti0.75Si0.11V0.03N
Ti0.80Si0.15N
Ti0.82V0.15N
TiN
2
36 37 38
Relative Intensity
TiN (200) TiN (220) TiN (311)
Ti0.52Si0.20V0.14N
Ti0.65Si0.11V0.15N
Ti0.75Si0.11V0.03N
Ti0.80Si0.15N
Ti0.82V0.15N
TiN
2
30 40 50 60 70 80
Table 3 shows the hardness, Young´s Modulus, residual stress and scratch test critical
load values of the coatings. V incorporation at the TiN and TiSiN systems slightly increased
the hardness and Young’s modulus value of coatings. Ti0.52Si0.20V0.14N coating exhibited the
lowest hardness and Young´s modulus as compared to the other films, a consequence of the
above referred loss of crystallinity. Residual stresses were similar for all coatings
(compressive; 3 – 4 GPa) suggesting that the shift of TiN peaks to higher angles should be
exclusively attributed to the presence of V and/or Si in solid solution. Moreover, by similar
reasons, the small enhancement in the hardness achieved with V additions is probably due to
solid solution hardening. Similar correlations could be drawn for TiSiN system in comparison
to TiN coating. In relation to the adhesion of films only two failure modes were observed on
the scratch tracks of the coatings: the first cracking (Lc1) and the first chipping (Lc2). Globally,
all the coatings are well adherent to the substrate exhibiting high Lc values with Ti0.82V0.15N
and Ti0.52Si0.20V0.14N being better and worse performing, respectively. In general, Si
incorporation in TiN system globally decreased the critical load values of the coatings.
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The wear rate of coatings tested at room temperature were dependent on the
counterpart material, as shown in Figure 2. Generally, when Al2O3 balls were used as
counterparts, the wear rate of coatings was much lower compared to sliding against HSS steel
balls. For reference coatings tested with Al2O3 balls, the tribological behaviour of TiN coating
was improved when vanadium was added; on the other hand, the addition of Si had a negative
effect. As it will be shown later, there is a correlation between the wear rates and the cracking
critical load values of these coatings, which suggests a possible relationship with the fracture
toughness of the coatings and its influence on the wear debris production.
When Si and V were added simultaneously, V had predominant influence. In fact, any
of the TiSiVN coatings showed lower wear rate than the TiN reference, with similar values to
Ti0.82V0.15N sample. It should be remarked that these coatings have, generally, similar Lc
values as Ti0.82V0.15N, i.e. higher Lc than the ones measured for Ti0.80Si0.15N . A similar trend
has been observed for other nitride coating systems containing V [7, 14]. Besides the
influence of V on the Lc values, the high wear resistance of V-rich coatings can also be
attributed to the formation of V lubricious oxides on the wear track of coatings (see further
discussion below). This can explain why Ti0.52Si0.20V0.14N shows lower wear rate than TiN
(and Ti0.80Si0.15N) in spite of its lower Lc value.
The wear rate of the coatings increased significantly when tested against HSS balls,
particularly for the coatings where the V-rich lubricious oxide was identified during the
contact with Al2O3 balls. The lower hardness of HSS ball associated with an easier
incrustation of hard wear debris in the sliding contact promoted abrasion of the wear track
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with increasing wear rates. Limited formation of vanadium oxides led to higher friction
coefficient, more energy dissipated and more oxide particles in the wear track, as it will be
discussed below.
As expected, the wear rate of the Al2O3 balls is much lower than that of softer HSS
balls. Different trends were found between the wear rates of the coatings and the balls: (i)
HSS ball wear rate increased when the coating wear decreased, and (ii) the wear rate of the
Al2O3 balls was very low and almost constant for all the tests. In both cases, the ball surfaces
were covered with the wear debris; however, HSS balls show large areas with shallow
scratches parallel to the movement of the ball and higher amount of adhered material.
14 Wear track
Al2O3 ball
Wear rate (10 mm /N.m)
12
HSS ball
10
3
8
-6
0
TiN Ti0.82V0.15N Ti0.80Si0.15N Ti0.65Si0.11V0.15N
a) Ti0.75Si0.11V0.03N Ti0.52Si0.20V0.14N
HSS ball
0.6
3
0.5
-6
0.4
0.3
0.2
0.1
0.0
TiN Ti0.82V0.15N Ti0.80Si0.15N Ti0.65Si0.11V0.15N
b) Ti0.75Si0.11V0.03N Ti0.52Si0.20V0.14N
Figure 2 - a) Wear rates of coatings tested against Al2O3 and HSS balls. b) Wear rates of Al2O3 and HSS
counterparts.
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Figure 3 a) and b) shows the friction coefficient of the coatings tested against Al 2O3
and HSS balls, respectively. As the radius of the wear tracks used for Al2O3 and HSS balls
tests were not the same and consequently, the sliding distances were different, the evolution
of the friction coefficient is plotted as a function of the total number of cycles (5000 cycles).
All friction curves exhibited two stages: running-in and steady state. Generally, the coatings
showed lower friction coefficient values against Al2O3 showing thus similar trend to the wear
rate. The highest friction coefficient was obtained for Ti0.80Si0.15N coating with values of 1.09
and 1.15 when tested against Al2O3 and HSS balls, respectively; it should be pointed out that
the wear rate was the highest as well. TiN coating presented also a high friction coefficient
when tested against Al2O3 ball, with a large variation in the values with the number of cycles.
Addition of Si and V had a significant impact on the frictional properties when Al2O3
ball was used as a counterpart. Compared to Ti0.80Si0.15N, the friction dropped to one half
when sufficient amount of vanadium was added and silicon content was kept low; the friction
of Ti0.65Si0.11V0.15N is below 0.5 being the lowest together with Ti0.82V0.15N coating. Testing
with HSS balls showed only limited influence of coating composition on the friction
coefficient. Again TiSiVN coatings exhibited the lowest friction, although it was only about
10% lower than that of Ti0.80Si0.15N and similar to that of TiN.
TiN Ti0.75Si0.11V0.03N
1.4 Ti0.82V0.15N Ti0.65Si0.11V0.15N
Ti0.80Si0.15N Ti0.52Si0.20V0.14N
1.2
Friction coefficient
1.0
0.8
0.6
0.4
0.2
0.0
0 1000 2000 3000 4000 5000
a) Number of cycles
Figure 3 - Friction coefficient vs number of cycles of coatings tested against: a) Al2O3 balls, b) HSS steel balls.
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TiN Ti0.75Si0.11V0.03N
1.4 Ti0.82V0.15N Ti0.65Si0.11V0.15N
Ti0.80Si0.15N Ti0.52Si0.20V0.14N
1.2
0.8
0.6
0.4
0.2
0.0
0 1000 2000 3000 4000 5000
b) Number of cycles
Figure 3 (continued).
The changes in the friction coefficient and wear rates can be only partially correlated
with the structural and mechanical properties of the films. To identify the dominant wear
mechanisms scanning electron microscopy with energy dispersive X-ray spectroscopy (SEM-
EDS) was used to analyze the worn surfaces and the wear debris produced in the contact.
Typical 2D profiles of the wear tracks of the coatings tested against Al2O3 are depicted in
Figure 4 a). Profilometer measurements showed that with the addition of V to TiN and
Ti0.80Si0.15N coatings the wear track became shallower. Ti0.80Si0.15N and TiN showed the
deepest and widest wear tracks among all the coatings, which gave rise to their highest wear
rates shown in Figure 2. The wear track of TiN coating was covered by fine scratches parallel
to the relative sliding movement suggesting abrasion as the dominant wear mechanism. The
scratches ended by the accumulation of the dragged material, which can well explain the
fluctuations observed in the friction curve (Figure 5)). EDS analysis performed in the wear
track revealed a mixture of as-deposited and oxidized coating material. Ti0.80Si0.15N coating
showed a similar wear mechanism as TiN; however, due to lower fracture toughness the worn
particles were larger and scratches deeper. The presence of large non-oxidized particles
originated from the coating led to an intensive abrasion of TiN and, particularly, Ti0.80Si0.15N
coating.
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8 8
Ti0.52Si0.20V0.14N Ti0.52Si0.20V0.14N
7 7
6 Ti0.65Si0.11V0.15N 6 Ti0.65Si0.11V0.15N
Distance (m)
Distance (m)
5 Ti0.75Si0.11V0.03N 5 Ti0.75Si0.11V0.03N
4
Ti0.80Si0.15N 4 Ti0.80Si0.15N
3 3
Ti0.82V0.15N Ti0.82V0.15N
2 2
TiN TiN
1 1
0 0
-0.4 -0.2 0.0 0.2 0.4 -0.4 -0.2 0.0 0.2 0.4
Distance (mm) Distance (mm)
a) b)
Figure 4 - 2D profiles of the wear track of coatings tested against: a) Al 2O3, b) HSS.
Scratches
Material accumulation
200 µm 20 µm
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oxide phases. It should be remarked that all the other zones of the wear track showed EDS
spectra identical to the as deposited films. After the test with Ti0.52Si0.20V0.14N coating, the
wear scar of the Al2O3 ball was covered with wear debris forming a transfer layer identified
by EDS analysis as a mixture of V-O, Ti-O and Si-O oxides. Friction induced oxidation of
wearing surfaces containing vanadium leading to formation of V-O, which acts as a solid
lubricant in sliding contact [2, 7, 8]. In this study, thick tribolayer consisting mostly of
vanadium oxide was indeed formed on the surface of Ti0.82V0.15N coating. However, addition
of silicon had detrimental effect on tribolayer formation and function during the sliding
process. It seems that Si somehow limits the formation of thick well-distributed tribolayer.
The higher wear rate and friction coefficient of Ti0.75Si0.11V0.03N in relation to
Ti0.65Si0.11V0.15N coating was in good accordance to the lower amount of “vermicular” debris
detected in the wear tracks. Although Ti0.82V0.15N and Ti0.65Si0.11V0.15N showed similar V
contents, the higher amount of transfer layer was formed in the case of Ti0.82V0.15N. The role
of Si is not clear, although it is possible that Si-O top layer hinders ion diffusion [15] and thus
protects the coating from oxidation required to form sufficient amount of lubricious vanadium
oxide.
The worn surface of Ti0.52Si0.20V0.14N film was found to be covered by fish-scale-like
agglomerates along the worn surface indicating severe shear deformation of the tribofilm
during the sliding test [8]. EDS analysis suggested the presence of V-O, Si-O and Ti-O oxides
in this zone. Additionally, EDS analysis showed significantly stronger peak of Si in relation to
the one found for the other TiSiVN films. Therefore, the oxide layer is dominated by Si-O.
Due to low fracture toughness large wear debris particles are detached during the running-in
stage; these particles are then either removed from the sliding contact (they were found on the
side of the wear track) or oxidized forming adhered oxide layer both on coating and ball
surface. However, due to the strong shear force that the ball imposes to the coating, and the
more brittle nature of Si-O-rich oxide layer, a stick-slip adhesion process will take place,
deforming and breakage the wear debris and giving rise to the above referred fish-like aspect.
Therefore, the tribological efficiency will be lower in this coating in relation to the other V-
containing films due to the combined effect of: i) lower fracture toughness, ii) less lubricious
V-O phase formed due the high Si content and iii) stick-slip adhesion wear mechanism.
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200 µm 200 µm
a) b)
Fish-like
Vermicular black wear debris
wear debris
5 µm 20 µm
c) d)
Figure 6 – SEM pictures of the wear track of: a) Ti0.65Si0.11V0.15N coating tested against Al2O3 ball, b)
Ti0.52Si0.20V0.14N coating tested against Al2O3 ball, c) and d) magnification of the wear tracks of Ti 0.65Si0.11V0.15N
and Ti0.52Si0.20V0.14N coatings, respectively.
Low hardness of HSS balls compared to Al2O3 resulted in high ball wear and wide
wear tracks (see Figure 4 b)). In fact, one might expect that hard coatings will resist much
more to sliding against relatively soft steel ball, whereas Al2O3 counterpart should induce
severe wear damage. However, Figure 2 and Figure 4 clearly indicate opposite behavior, i.e.
higher wear with HSS balls.
Reference TiN and Ti0.80Si0.15N coatings exhibited slightly different wear track
appearance. Figure 7 shows the typical morphology of the worn surfaces of these coatings and
the corresponding wear scars of the balls. EDS analysis performed on the transfer layer of
TiN coating, which was preferential formed in the central part of the wear track, evidenced
material transfer from the ball to the coating surface. The transfer layer thus consisted of
titanium and iron oxides. However, no oxygen was detected by EDS outside of the center part
of unworn coating. Although the wear scars of the counterparts were covered with as-
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deposited and oxidized material from the coating for both TiN and Ti0.80Si0.15N coatings, the
latter exhibited higher material transfer to the HSS ball and, particularly, larger wear debris
particles. Hard non-oxidized particles embedded into compact adhered tribolayer act as sharp
cutting edges causing abrasive damage and accelerating wear. On the other hand, relatively
thick and well-adhered layer on the ball wear scar protected ball material from further
damage; thicker layer built during sliding against on Ti0.80Si0.15N resulted in lower ball wear
(Figure 2 b)). It is clear here that higher amount of non-oxidized debris adhered on the ball
protects the ball itself but produces higher wear of the coating; the absence of well-developed
oxide layer on the coating surface is detrimental decreasing coating wear resistance.
Transfer
layer
200 µm 200 µm
a) b)
400 µm 400 µm
c) d)
Figure 7 – Wear tracks induced by HSS ball sliding in: a) TiN coating, b) Ti0.80Si0.15N coating. Wear scar on
HSS balls tested against: c) TiN coating, d) Ti0.80Si0.15N coating.
Absence of continuous tribolayer in the wear track was observed as well for the
coatings containing vanadium. Typical worn surfaces of the coatings (Ti 0.82V0.15N and
Ti0.65Si0.11V0.15N) shown in Figure 8 were covered by stripes of adhered material, which
covered only small part of the wear tracks, see Figures 8 a) and b), indicating that the wear is
driven by a combination of adhesion and polishing wear. EDS analysis revealed that the
stripes were composed of fully oxidized (i.e. no nitrogen observed) coating elements mixed
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with iron originating from the ball. Outside of the stripes no oxygen was identified by EDS.
The phenomenon of adhesive wear behaviour was believed to be a result of hard film and
relatively soft bearing ball [16]. Nevertheless, some of the V-O oxide present on the stripe
zones should still provide a lubricious effect since a decrease of the wear rate with V
incorporation is observed. This behaviour is enhanced in case of Ti0.52Si0.20V0.14N coating
where an even higher wear rate is observed; its lower fracture toughness, due to the higher
silicon content, constrains the V-O phase formation decreasing even more its content on the
wear track. Ti0.82V0.15N film displayed the lowest wear rate of all the coatings due to the
higher amount of V-O oxide formed at the surface.
In summary, we demonstrated that V additions successfully reduced the wear rate and
friction coefficient of TiSiN system; moreover we also showed that the tribological behaviour
of films is strongly dependent on the counterpart material due to the presence of dissimilar
wear mechanisms.
200 µm b) 200 µm
a)
20 µm 20 µm
c) d)
Figure 8 - SEM pictures of the wear track tested against HSS balls of: a) Ti0.82V0.15N coating, b) Ti0.65Si0.11V0.15N
coating. c) and d) magnification of the wear tracks of Ti 0.82V0.15N and Ti0.65Si0.11V0.15N coatings, respectively.
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4. Conclusion
The results showed that alloying TiSiN system with increasing V content improves the
resistance to wear and decreases the friction coefficient; however, the tribological behaviour
is strongly dependent of the counterbody ball material. SEM and Raman analysis revealed
that the influence of V was due to the formation of V-O lubricious oxide on the contact zone.
When Al2O3 balls were used in the tests, the wear was driven by polishing wear and oxides
containing V-O are spread all over the wear track. On the other hand, when tested against
HSS balls, adhesion wear took place with material transfer from the ball to specific zones of
the wear track, being the lubricious effect of V-O significantly reduced. Ti0.80Si0.15N film
displayed the lowest wear resistance among all the tested coatings due to its lower fracture
toughness which promotes the formation of bigger wear debris acting as strong abrasion
players.
Acknowledgments
This research is sponsored by FEDER funds through the program COMPETE – Programa
Operacional Factores de Competitividade – and by national funds through FCT – Fundação
para a Ciência e a Tecnologia –, under the project: PTDC/EME-TME/122116/2010, as well as
the grant (SFRH/BD/68740/2010).
References
[3] E. Lugscheider, O. Knotek, K. Bobzin, S. Bärwulf, Surf. Coat. Technol., 133–134 (2000)
362-368.
[4] N. Fateh, G.A. Fontalvo, G. Gassner, C. Mitterer, Wear, 262 (2007) 1152-1158.
204
Filipe Daniel Fernandes
Annex I
[6] Y. Qiu, S. Zhang, J.-W. Lee, B. Li, Y. Wang, D. Zhao, Applied Surface Science, 279
(2013) 189-196.
[9] K. Kutschej, P.H. Mayrhofer, M. Kathrein, P. Polcik, C. Mitterer, Surf. Coat. Technol.,
200 (2005) 1731-1737.
[10] F. Fernandes, A. Loureiro, T. Polcar, A. Cavaleiro, Applied Surface Science, 289 (2014)
114-123.
[11] European Committee for standardization, European Stardard DIN ENV 1071-3 (1994),
draft European Standard prEN 1071-3, (2002).
[12] G.G. Stoney, Proceedings of the Royal Society of London. Series A, 82 (1909) 172-175.
[13] S.H. Kim, J.K. Kim, K.H. Kim, Thin Solid Films, 420–421 (2002) 360-365.
[14] K. Kutschej, P.H. Mayrhofer, M. Kathrein, P. Polcik, C. Mitterer, Surf. Coat. Technol.,
188–189 (2004) 358-363.
[16] C.-L. Chang, C.-T. Lin, P.-C. Tsai, W.-Y. Ho, D.-Y. Wang, Thin Solid Films, 516 (2008)
5324-5329.
205
Filipe Daniel Fernandes