VSEBINA – CONTENTS
PREGLEDNI ^LANKI – REVIEW ARTICLES
Aluminium: the metal of choice
Aluminij: kovina izbire
M. J. F. Gándara . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 261
Nanotechnology for ballistic materials: from concepts to products
Nanotehnologija za balisti~ne materiale: od izhodi{~ do proizvoda
V. M. Castaño, R. Rodríguez . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 267
Multi-functional KIc-test specimen for the assessment of different tool- and high-speed-steel properties
Ve~funkcijski KIc-preizku{anec za dolo~anje razli~nih lastnosti orodnih in hitroreznih jekel
V. Leskov{ek, B. Podgornik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 273
Nanosilica-reinforced polymer composites
Polimerni kompoziti oja~ani z nanosiliko
M. Conradi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 285
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
The fatigue behaviour of aluminium foam
Vedenje aluminijevih pen pri preizkusu utrujenosti
M. Nosko, F. Siman~ík, R. Florek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295
Addition of strontium to an Mg-3Sn alloy and an investigation of its properties
Dodatek stroncija zlitini Mg-3Sn in preiskava njenih lastnosti
M. Öbekcan, A. Ayday, H. ªevýk, S. C. Kurnaz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 299
The effect of binder on chemically precipitated hydroxyapatite during spray drying
Vpliv veziva na kemijsko izlo~eni hidroksiapatit med atomizacijo
F. E. Baþtan, E. Demiralp, Y. Y. Özbek, F. Üstel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 303
Plasma electrolytic saturation of 316 L stainless steel in an aqueous electrolyte containing urea and ammonium nitrate
Plazemsko elektrolitsko nasi~enje nerjavnega jekla 316 L v vodnem elektrolitu s se~nino in amonijevim nitratom
L. C. Kumruoglu, A. Özel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 307
Minimization of surface defects by increasing the surface temperature during the straightening of a continuously cast slab
Zmanj{evanje povr{inskih napak z zvi{anjem temperature povr{ine kontinuirno ulitega slaba med ravnanjem
J. Stetina, T. Mauder, L. Klimes, F. Kavicka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 311
Influence of process parameters on the corrosion resistance of corrugated austenitic and duplex stainless steels
Vpliv procesnih parametrov na korozijsko odpornost rebrastih avstenitnih in dupleksnih nerjavnih jekel
S. M. Alvarez, A. Bautista, F. Velasco . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 317
Concentration and path-length dependence on the Faraday rotation of magnetic fluids based on highly water-soluble
Fe3O4/PAA nanoparticles synthesized by a high-temperature hydrolysis method
Odvisnost koncentracije in dol`ine poti od Faradayevega vrtenja magnetnih teko~in na osnovi visokovodotopnih nanodelcev Fe3O4/PAA,
sintetiziranih z metodo visokotemperaturne hidrolize
S. Küçükdermenci, D. Kutluay, E. Çelik, Ö. Mermer, Ý. Avgýn . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 323
The possibility of copper corrosion protection in acidic media using a thiazole derivative
Mo`nost protikorozijske za{~ite bakra v kislem mediju z uporabo derivatov tiazola
\. Va{tag, S. Apostolov, M. Had`istevi}, M. Sekuli} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 329
Energy- and time-saving low-temperature thermomechanical treatment of low-carbon plain steel
Prihranki energije in ~asa pri nizkotemperaturni termomehanski obdelavi malooglji~nega plo{~atega jekla
H. Jirkova, D. Hauserova, L. Kucerova, B. Masek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 335
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 47(3)259–400(2013)
MATER.
TEHNOL.
LETNIK
VOLUME 47
[TEV.
NO. 3
STR.
P. 259–400
LJUBLJANA
SLOVENIJA
MAY–JUNE
2013
Electrochemical characterization of the nano Py/DDS/SiO2 film on a copper electrode
Elektrokemijska karakterizacija nanoplasti Py/DDS/SiO2 na bakreni elektrodi
M. Sharifirad, F. Kiani, F. Koohyar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 341
Sewage-sludge stabilization with biomass ash
Stabiliziranje komunalnega mulja s pepelom biomase
P. Pav{i~, D. O{tir, A. Mladenovi~, S. Kramar, M. Dolenec, P. Bukovec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 349
Dredged mud from the port of Koper – civil engineering applications
Mulj iz luke Koper – uporabnost v gradbeni{tvu
A. Mladenovi}, @. Poga~nik, R. Mila~i~, A. Petkov{ek, F. Cepak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 353
Automated diagnostics of damage to an aluminum alloy under the conditions of high-cycle fatigue
Avtomatizirana diagnostika po{kodbe aluminijeve zlitine pri visoko-cikli~nem utrujanju
P. Maruschak, I. Konovalenko, M. Karuskevich, V. Gliha, T. Vuherer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 357
Welded aluminium and magnesium alloys – corrosion and mechanical properties for refrigeration compressors in comparison
with deep-drawing steel
Varjene aluminijeve in magnezijeve zlitine – korozijske in mehanske lastnosti za kompresorje hladilnikov v primerjavi z jeklom za globoko
vle~enje
K. G. Kerschbaumer, R. Vallant, N. Enzinger, C. Sommitsch . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 363
Water-vapour plasma treatment of cotton and polyester fibres
Obdelava bomba`nih in poliestrskih vlaken s plazmo vodne pare
J. Vasiljevi}, M. Gorjanc, R. Zaplotnik, A. Vesel, M. Mozeti~, B. Simon~i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 379
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Influence of sample direction on the impact toughness of the API-X42 microalloyed steel with a banded structure
Vpliv usmerjenosti vzorcev na udarno `ilavost mikrolegiranega jekla API-X42 s trakavo mikrostrukturo
A. Salimi, H. Monajati Zadeh, M. R. Toroghinejad, D. Asefi, A. Ansaripour . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 385
Experimental and numerical investigation of an air-PCM heat-storage unit
Eksperimentalna in numeri~na preiskava enote zrak – PCM za shranjevanje toplote
T. Mauder, P. Charvat, M. Ostry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 391
Experimental investigation of a heat-transfer coefficient
Preiskave koeficienta prenosa toplote
M. Chabi~ovský, M. Raudenský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 395
IN MEMORIAM
Prof. dr. Jo`e Rodi~. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 399
M. J. FREIRÍA GÁNDARA: ALUMINIUM: THE METAL OF CHOICE
ALUMINIUM: THE METAL OF CHOICE
ALUMINIJ: KOVINA IZBIRE
María Josefa Freiría Gándara
Xunta de Galicia, Consellería de Educación e Ordenación Universitaria, 15704 Santiago de Compostela, La Coruña, Spain
josefa.freiria@yahoo.es
Prejem rokopisa – received: 2012-08-21; sprejem za objavo – accepted for publication: 2012-11-21
This article summarizes the importance of aluminium as the metal of choice for many applications. Aluminium is a lightweight,
durable metal. It is silvery in appearance when freshly cut, is a good conductor of heat and electricity, and is easily shaped by
moulding and extruding. Aluminium has two main advantages when compared with other metals. Firstly, it has a low density,
about one-third that of iron and copper. Secondly, although it reacts rapidly with the oxygen in air, it forms a thin, tough and
impervious oxide layer that resists further oxidation. This removes the need for surface-protection coatings such as those
required with other metals, in particular with iron.
All the indications are that the growth in the use of aluminium will likely accelerate. It is expected that in the near future the use
of aluminium with specifically improved properties will grow in many applications, meeting the increased economic and
ecological demands.
Considering the entire life-cycle of an automobile, from the extraction of materials to the final disposal, including recycling and
reuse applications, aluminium proves to be a potential alternative to steels in future automotive applications.
Keywords: corrosion resistance, aluminium alloys, aluminium applications, aluminium in vehicles, CO2 emission, aluminium
recycling
^lanek povzema pomembnost aluminija kot izbrane kovine z mnogimi mo`nostmi uporabe. Aluminij je lahka, zdr`ljiva kovina.
Je srebrne barve, ko je sve`e odrezan, je dober prevodnik toplote in elektrike in se lahko oblikuje z ulivanjem v kokile in z
ekstruzijo. Aluminij ima v primerjavi z drugimi kovinami dve glavni prednosti. Prva je majhna gostota, okrog ene tretjine
gostote `eleza ali bakra. Druga: ~eprav hitro reagira s kisikom iz zraka, tvori tanko, trdo in neprepustno oksidno plast, ki
prepre~uje nadaljnjo oksidacijo. Zato ni potrebe po povr{inski za{~iti, kot jo zahtevajo druge kovine, {e posebno `elezo.
Ka`e, da se bo rast porabe aluminija {e nadaljevala. V bli`nji prihodnosti se pri~akuje, da bo zaradi pove~anja ekonomskih in
ekolo{kih zahtev uporaba aluminija s posebnimi, izbolj{animi lastnostmi rastla na mnogih podro~jih uporabe.
Z upo{tevanjem celotnega trajnostnega cikla avtomobila, od pridobivanja materiala do odlaganja, vklju~no z recikliranjem in
mo`nostjo nadaljnje uporabe, aluminij dokazuje, da je mogo~a alternativa jeklom v prihodnji proizvodnji avtomobilov.
Klju~ne besede: odpornost proti koroziji, aluminijeve zlitine, uporaba aluminija, aluminij v vozilih, emisija CO2, recikliranje
aluminija
1 INTRODUCTION
Aluminium is a light, conductive, corrosion-resistant
metal with a strong affinity for oxygen. This combina-
tion of properties has made it a widely used material,
with applications in the aerospace, architectural con-
struction and marine industries, as well as many dome-
stic uses.
Aluminium is the second most widely used metal in
the world today. It is used extensively in aircraft, in
building construction, and in consumer durables such as
fridges, cooking utensils and air conditioners, as well as
in food-processing equipment and cans.
Aluminium is not found in its metallic form in
nature. It occurs as bauxite, a mixture of aluminium
oxides, iron oxides and clay. Manufacturing aluminium
metal from bauxite is a complex process.
Aluminium is one of the most important metals used
by modern societies. Aluminium´s combination of physi-
cal properties results in its use in a wide variety of pro-
ducts, many of which are indispensable in modern life.
Because of its light weight and electrical conductivity,
aluminium wire is used for the long-distance transmis-
sion of electricity. Aluminium’s strength, light weight,
and workability have led to increased use in transpor-
tation systems, including light vehicles, railcars, and
aircraft in efforts to reduce fuel consumption. Alumi-
nium’s excellent thermal properties and resistance to
corrosion have led to its use in air conditioning, refrige-
ration, and heat-exchange systems. Finally, its malleabi-
lity has allowed it to be rolled and formed into very thin
sheets used in a variety of packaging.
In 1903, the Wright brothers made aviation history
when they achieved the world’s first flight powered by a
lightweight engine made with aluminium components.
Today, aluminium is fundamental to the aviation indu-
stry. It was in the 1920s that aluminium shipping appli-
cations started to expand due to new alloys becoming
available for marine applications. In the 1980s, alumi-
nium emerged as the metal of choice for reducing
running costs and improving the acceleration of metros,
tramways, intercity and high-speed trains. The average
volume of aluminium used in passenger cars was
significantly increased in 2000.
The examples given for its success prove the major
breakthroughs in automotive applications for aluminium
that have been achieved during recent years by develo-
ping innovative light-weight and cost-efficient solutions.
Materiali in tehnologije / Materials and technology 47 (2013) 3, 261–265 261
UDK 669.71:669.715 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 47(3)261(2013)
Vehicle manufacturers must constantly improve their
performance at minimum costs. The choice of a material
will therefore depend on its price, its mechanical
properties and its impact on vehicle production costs.
Due to its low weight, good formability and corrosion
resistance, aluminium is the material of choice for many
automotive applications, such as the chassis, auto body
and many structural components.1–4
2 ALUMINIUM PRODUCTION
The major raw materials required for aluminium pro-
duction are alumina, carbon, power, aluminium fluoride
and cryolite. The aluminium industry consumes nearly
90 % of the bauxite mined; the remainder is used in
abrasives, cement, ceramics, chemicals, metallurgical
flux, refractory products, and miscellaneous products.
Virtually all the alumina commercially produced from
bauxite is obtained using a process patented by Karl
Josef Bayer (Austria) in 1888.
Primary aluminium is produced by the electrolysis of
alumina dissolved in molten fluoride salt. The process
was independently invented in 1886 by Charles Martin
Hall (United States) and Paul Louis Toussaint Héroult
(France) and underwent continual improvement over the
years. The electrolysis of alumina to produce aluminium
involves the use of aluminium fluoride, carbon anodes,
and large amounts of electricity.5–7
Taking the purity grades of aluminium into account,
the aluminium content is usually the main consideration
and other elements are considered only as impurities.
The common purity grades of aluminium are listed in
Table 1.8,9
3 CORROSION RESISTANCE
Aluminium’s well-known corrosion resistance is an
obvious advantage in road transport. It contributes to a
long service life, especially in vehicles that work in
conditions that can cause serious corrosion problems.
Usually, no painting or other surface protection is
required and it is easy to clean. Maintenance is therefore
kept to a minimum.
Corrosion is an electrochemical interaction between a
metal and its environment which results in changes to the
properties of the metal and which may often lead to
impairment of the function of the metal, the environ-
ment, or the technical system of which these form a part.
Corrosion can occur locally (pitting), or it can extend
across a wide area to produce a general deterioration.
A clean aluminium surface is very reactive and will
react spontaneously with air or water to form aluminium
oxide. This oxide builds a natural protective layer on
each aluminium surface with a thickness of around 1–10
nm. The oxide layer is chemically very stable, has a
good adhesion to the metal surface, repairs itself and
protects the aluminium from further corrosion. The oxide
layer can be destroyed in strong acidic or alkaline
environments or where aggressive ions are present.
Aggressive ions can destroy the layer locally and lead to
local corrosion attack (pitting). A typical case for this
reaction is the contact between aluminium and chloride
ions, which are present in seawater or road salts. Some
alloying elements might increase the corrosion resistance
of the oxide layer, while others can weaken it. Vehicle
manufacturers or fleet operators should contact the alu-
minium supplier in any case of critical working condi-
tions, like elevated temperatures or aggressive loads.
Some general rules need to be applied to prevent
corrosion, in most cases to prevent any kind of water trap
or areas where condensation can occur.
Although aluminium can be used without any surface
protection and keeps its natural beauty throughout its
life, it is most likely to use different surface treatment
methods to optimise its attractiveness and optical
appearance and to protect it from severe atmospheric
conditions.10
4 ALUMINIUM ALLOYS
In aluminium alloys other elements are deliberately
added to improve the properties in some way. Many
alloys have been developed, the aim being to improve the
strength while retaining the desirable properties of
aluminium, most notably its lightness and corrosion
resistance. In general though, while the addition of an
alloying element increases the strength, it reduces the
resistance to corrosion, making a compromise of the
properties necessary. A possible exception to this is
magnesium alloys, which have improved corrosion
resistance in marine environments. Aluminium-copper
alloys have very poor resistance to corrosion, and sheets
are often produced in sandwich form with thin layers of
pure corrosion-resistant aluminium on the outside. A
summary of typical alloys is given in Table 2.6,7,9
M. J. FREIRÍA GÁNDARA: ALUMINIUM: THE METAL OF CHOICE
262 Materiali in tehnologije / Materials and technology 47 (2013) 3, 261–265
Table 1: Purity grades of aluminium
Tabela 1: Stopnje ~istote aluminija
Aluminium content in
mass fractions, w/%
Major impurities in mass fractions, w/%
Some typical uses
Silicon Iron
99.95 (high purity) <0.006 <0.006 Extrusion joinery, electrical conductor, anodic trim, foil
99.80 <0.15 <0.15 Plumbing, reflectors, jewellery
99.50 <0.25 <0.40 Chemical plant, tanks, tubes
99.50 w(Si + Fe) < 1.0 % Pots, pans, sheet-metal work
Table 2: Aluminium alloys
Tabela 2: Zlitine aluminija
Major alloy
element
Content in
mass fractions,
w/%
Product Some typicaluses
Copper Up to 4.5
Sheet
Extrusions
Castings
High-strength
aircraft parts
Manganese 1.25 Sheet
Sheet-metal
work, pots,
pans, etc.
Silicon Up to 13 Castings
Motor parts,
castings of all
types
Magnesium
and Silicon
0.7 Mg
0.4 Si
Sheet
Extrusions
Architectural
extrusions
Magnesium Up to 5 Sheet
Marine uses,
boats, fish
boxes, beer-
can lids, etc.
Zinc,
Magnesium
and Copper
5.8 Zn
2.5 Mg
1.4 Cu
Sheet
Extrusions
High-strength
aircraft
Aluminium alloys used in the manufacture of
commercial vehicles and their accessories are easy to
process. They lend themselves to a variety of shaping
and joining techniques. Correctly used, aluminium alloys
have been developed to offer optimum corrosion resi-
stance in all environments. Just one example: the
widespread use of unpainted aluminium in marine
applications.
Aluminium alloys tailored by suitable variations in
chemical composition and processing best fit many
requirements, like the non-heat-treatable Al-Mg alloys
used in chassis optimized for superb resistance against
inter-crystalline corrosion and concurrent high strength
or the heat-treatable Al-Mg-Si alloys for extrusions and
auto-body sheet modified for improved age-hardening
response during the automotive paint bake cycle.
Moreover, Al-Mg-Mn alloys show an optimum
combination of formability and strength achieved by the
mechanism of solid solution and deformation hardening
due to their specific high-strain hardening. Further
improvements in the properties required for specific
applications (e.g., surface appearance, corrosion resi-
stance, thermal stability) have been achieved by small
additions of other alloy elements and/or modified pro-
cessing routes, e.g., stretcher strain free (SSF) sheet,
avoiding Lüders-lines. Non-heat-treatable Al-Mg-Mn
alloys are often applied for automobile parts in larger
quantities as hot- and cold-rolled sheet and hydroformed
tubes, due to their good formability, which can always be
regained during complex forming operations by inter-
annealing, where quenching is needed for age hardening.
In chassis parts or wheel applications the benefit is
twofold, since the weight reduction in the unsprung mass
of moving parts additionally enhances driving comfort
and reduces noise levels.11–19
5 ALUMINIUM APPLICATIONS: ALUMINIUM
IN VEHICLES
Aluminium in its pure form is a very soft metal and
hence not suited for structural applications. Thanks to
the addition of alloying elements such as copper, manga-
nese, magnesium, zinc, etc., and thanks to adequate
production processes, the physical and mechanical
properties can be varied over a great range, making it
possible to have suitable alloys for literally all appli-
cations.
There are challenges with aluminium. Perhaps the
biggest challenge is the history of the auto industry. The
many years of experience based on steel technology
represents a significant barrier for aluminium, especially
in the areas of manufacturing, i.e., forming and joining,
which are critical to the automotive industry.
In the eighteenth century, aluminium was very expen-
sive, in spite of the fact that aluminium is a very
abundant metal in the earth’s crust. Since that time the
cost of aluminium has been on a continuous and steep
decline in price based on technological advancements.
This trend will certainly continue since there are a great
number of initiatives, which will continue this decline in
prices.
Aluminium can be used for the car body structure
and there can be a weight advantage of at least 30 %
without any loss of performance. In some cases where
very high strengths are demanded, 7xxx series alloys can
be used to maintain this significant weight advantage.
For a large volume, aluminium solutions are the most
cost effective. Castings will be applied for areas where a
strong part integration is feasible. Extrusions can easily
be applied as straight profiles, but also forming of an
extruded profile is a competitive process for high
volumes, e.g., as bumper beams. Aluminium is the ideal
light-weight material as it allows a weight saving of up
to 50 % over competing materials in most applications
without compromising safety.
Due to the positive experiences and from former
successful applications its volume fraction used in cars
of all classes and all sizes will grow significantly.
Applying full knowledge about the physical processes
involved and the microstructure/properties correlation a
tuning of properties is possible in order to produce
optimum and stable products required for the high
demands in automobile applications.
The automotive industry has more than doubled the
average amount of aluminium used in passenger cars
during recent years and will do even more so in the
coming years. The automotive industry, in close
co-operation with the aluminium industry, has developed
and introduced numerous innovative light-weight solu-
tions based on established and improved aluminium
alloys and optimized aluminium-oriented car design.
Synergic effects together with a multi-material exploita-
tion can guarantee an optimum design solution. One of
M. J. FREIRÍA GÁNDARA: ALUMINIUM: THE METAL OF CHOICE
Materiali in tehnologije / Materials and technology 47 (2013) 3, 261–265 263
the main advances of aluminium is its availability in a
large variety of semi-finished forms, such as shape
castings, extrusions and sheet, all suitable for mass
production and innovative solutions. Compact and highly
integrated parts meet the strong demands for high
performance, quality and cost-efficient manufactura-
bility. Challenges involved here are mainly joining and
surface treatment issues for which many suitable
solutions have been developed. Aluminium semis are
applied as castings, extrusions and sheet increases, e.g.,
in engine blocks and power-train parts, space frames,
sheet structures or as closures and hang-on parts and
other structural components.
Material-specific processing routes and individual
solutions have been developed in close cooperation with
partners and suppliers. With a sound knowledge about
the specific material properties and effects excellent
light-weight solutions for automotive applications have
been successfully applied by the automobile industries.
Intensive and continuous collaboration of material
suppliers and application engineers provided optimum
solutions for sometimes contradicting aspects of the
specific requirements, e.g., for the specific material
selection and optimum combinations of strength and
formability. Safety has become a crucial issue for vehicle
manufacturers. Aluminium has the advantage of being
much stronger than steel on a weight basis, so that with
proper design a lighter aluminium vehicle can be
expected to protect vehicle passengers as well (or better)
than a heavier steel vehicle.
The manufacturing flexibility of aluminium also
represents a real advantage for the metal in auto applica-
tions. Thus, aluminium is easier to extrude and cast than
steel. Other processes, such as semi-solid forming and
forging, are finding niches in the automotive business as
well. Aluminium sheet applications are increasing at a
rapid rate in structural, exterior panels and closure
panels.
The auto industry is highly competitive and increas-
ingly global. Automakers are being challenged not only
to meet the expectations of shareholders and customers,
but also to answer the growing environmental concerns
of society. Aluminium is the third most highly used
material in vehicles and is rapidly gaining on its rival
materials (iron and steel). Aluminium contributes to the
reduction of CO2 emissions from road transport. In
recent years the potential problem of global warming has
provided additional pressure on automakers worldwide
to improve fuel economy. Carbon dioxide is considered
to be the biggest anthropogenic contribution to global
warming. It results primarily from the combustion of
fossil fuels, hence the pressure on autos to burn less fuel.
However, when one considers material substitutions,
such as aluminium for steel, it is necessary to take
account of the entire process of obtaining, processing,
using and recycling the material.20,21
The automotive industry is known worldwide as
being technically advanced and innovative. Based on
economic and political pressure to reduce fuel consump-
tion and CO2 emissions the efforts for light weight in
automobile design and constructions have increased
significantly and specific solutions based on the inten-
sive use of aluminium as modified or new alloys have
been developed in recent decades.
Reducing manufacturing costs and tailpipe emissions
by using light-weight materials which can be easily
recycled or reused are among the major issues in today’s
automobile industry. The growing emphasis on total cost
and environmental impact has forced the life-cycle cost
issue to be the driving factor. Weight savings in the
overall car mass is considered to be a major research
focus. Aluminium is proven to be among the potential
materials capable of achieving weight reduction without
sacrificing either vehicle safety or performance. Despite
significant technological advantages in aluminium
alloys, the use of aluminium alloys to replace traditional
materials such as steels has been slow due to the lack of
a comprehensive economic analysis of the entire life-
cycle of automotive products. In considering the total
life-cycle of an automobile covering four stages
(pre-manufacturing, manufacturing, use, and post-use), it
is apparent that during the operational (use) stage of a
vehicle, aluminium is proven to be a reliable alternative
for traditional materials currently used in automotive
body structures, largely due to its cost-effectiveness and
superior performance due to light weight. With gas price
variations, the initial cost advantage of using steel in
body components gained in pre-manufacturing and
manufacturing stages can be overcome during the opera-
tional (use) stage of the vehicle, since the lighter
alternative provides significant savings in terms of fuel
consumption and consequently the generation of air-
borne gas emissions. Also, the superior recyclability and
reusability of aluminium in the post-use stage outweighs
that of traditional materials, despite the higher costs
involved in producing primary aluminium.
Knowing that the greatest opportunity for weight
savings comes from the body structure and exterior
closure panels, and that additional weight reduction can
be achieved by downsizing the other components such as
engine components, it can be considered as achieving a
weight reduction by replacing the conventional material
used in the vehicle’s construction (i.e., steel) with a
lighter mass equivalent material (i.e., aluminium), main-
taining the same vehicle design and using the same
manufacturing processes for the body components.22–28
6 ALUMINIUM RECYCLING
Aluminium is easily and economically recycled.
Aluminium does not degrade during the recycling pro-
cess, which means it can be recycled over and over
again. Recycling aluminium reduces the need for raw
M. J. FREIRÍA GÁNDARA: ALUMINIUM: THE METAL OF CHOICE
264 Materiali in tehnologije / Materials and technology 47 (2013) 3, 261–265
materials and reduces the use of valuable energy
resources. Moreover, recycling reduces the amount of
waste in landfill.
Aluminium is a valuable material to recycle due to
the large amount of energy and resources used in the
initial manufacture. It can be infinitely recycled.
Recycled aluminium is made into aircraft, automobiles,
bicycles, boats, computers, cookware, gutters, sidings,
wires, cans, etc.
7 REFERENCES
1 M. J. Crooks, R. E. Miner, Journal of Metals, 48 (1996) 7, 13–15
2 Y. Kurihara, Journal of Metals, 45 (1993) 11, 32–33
3 G. P. Syukla, D. B. Goel, P. C. Pandey, All India Seminar on Alumi-
nium, New Delhi, 1978
4 Report of Investigation on the Technical Trend of Patent Applica-
tions in 2004: Automobile Weight reduction Technologies, Japan
Patent Office, 2005
5 J. E. Gould, Projection welding, ASM Handbook, ASM Internatio-
nal, Metals Park, Ohio, 6, 1993, 230–237
6 A. S. Russell, Aluminum, McGraw-Hill Encyclopedia of Science &
Technology, McGraw-Hill, New York 1997
7 D. Altenpohl, Aluminum Viewed from within: An Introduction into
the Metallurgy of Aluminum Fabrication, Aluminium-Verlag,
Düsseldorf 1982
8 U. M. J. Boin, M. Bertram, Journal of Metals, 5 (2005), 26–33
9 K. Grjotheim, B. J. Welch, Aluminium Smelter Technology: A Pure
and Applied Approach, 2nd ed., Aluminium-Verlag, 1988
10 G. D. Davis, J. S. Ahearn, J. D. Venables, J. Vac. Sci. Technol., A2
(1984), 763–766
11 I. N. Fridlyander, V. G. Sister, O. V. Grushko, V. V. Berstenev, L. M.
Sheveleva, L. A. Ivanova, Metal Science and Heat Treatment, Sep-
tember (2002), 3–9
12 M. P. Meinel, Product Engineering, 23 (1951), 163–165
13 T. R. Gauthier, W. J. Rowe, Mechanical Engineering, 70 (1948),
505–514
14 W. B. F. Mackay, R. L. Dowdell, Metal Progress, 56 (1949),
331–336, 404, 406
15 P. O’Keefe, Materials and Methods, 33 (1951), 90–104
16 T. P. Fisher, British Foundryman, 66 (1973), 71–83
17 R. F. Shyu, C. T. Ho, Journals of Materials Processing Technology,
171 (2006) 3, 411–416
18 H. Xu, Q. Han, T. T. Meek, Materials Science and Engineering: A,
473 (2008) 1–2, 96–104
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(2009) 5, 1604–1611
20 A. A. Fernandez-Bremauntz, M. R. Ashmore, Atmospheric environ-
ment, 29 (1995), 525–532
21 P. G. Flachsbart, Journal of exposure analysis and environmental
epidemiology, 5 (1995), 473–495
22 R. V. Vanden Berg, Product Engineering, 22 (1951), 179–186
23 W. D. Stewart, Mechanical Engineering, 75 (1953), 450–455
24 F. H. Froes, Journals of Materials & Design, 10 (1989) 3, 110–120
25 J. Jennings, J. E. Gould, Welding Journal, 87 (2008) 10, 26–30
26 Energy consumption Trend in Transport Sector: 2-2-1. Analysis on
private passenger cars’ contribution to the total energy consumption.
Home page provided by the Energy Conservation Center, Japan,
http://www.eccj.or.jp/transportation/2-1-1.html
27 M. Surappa, Sadhana, 28 (2003) 1, 319–334
28 U.S. Department of Energy (DOE), Aluminum Project Fact Sheet:
Dynamic Inert Metal Anodes, 2001 Office of Industrial Technolo-
gies, Washington D.C., 2001
M. J. FREIRÍA GÁNDARA: ALUMINIUM: THE METAL OF CHOICE
Materiali in tehnologije / Materials and technology 47 (2013) 3, 261–265 265

V. M. CASTAÑO, R. RODRÍGUEZ: NANOTECHNOLOGY FOR BALLISTIC MATERIALS ...
NANOTECHNOLOGY FOR BALLISTIC MATERIALS:
FROM CONCEPTS TO PRODUCTS
NANOTEHNOLOGIJA ZA BALISTI^NE MATERIALE:
OD IZHODI[^ DO PROIZVODA
Víctor M. Castaño, Rogelio Rodríguez
Centro de Física Aplicada y Tecnología Avanzada, Universidad Nacional Autónoma de México, Boulevard Juriquilla 3001, Santiago de
Querétaro, 76230 Querétaro, México
castano@fata.unam.mx
Prejem rokopisa – received: 2012-08-31; sprejem za objavo – accepted for publication: 2012-11-16
The main trends and materials in protection technologies are briefly reviewed, emphasizing the properties and limitations of
p-aramid fibres, widely used in armour systems, particularly in terms of their susceptibility to UV radiation, humidity and
chemical attacks. Then, a novel nanotechnology capable of effectively diminishing these effects is described, as well as its
application for an actual commercial ballistic vest.
Keywords: nanotechnology, ballistic materials, aramid fibres, bullet-proof materials, nanoparticles, degradation
Podan je pregled glavnih usmeritev in materialov v tehnologijah za{~ite s poudarkom na lastnostih in omejitvah p-aramidnih
vlaken, ki se uporabljajo v oboro`itvenih sistemih s poudarkom na njihovo ob~utljivost za UV-sevanje, vlago in kemi~ni napad.
Nato so opisane nove nanotehnologije, ki so sposobne zmanj{ati te vplive, kot tudi njihova uporaba pri dejanskem komercial-
nem balisti~nem telovniku.
Klju~ne besede: nanotehnologija, balisti~ni materiali, aramidna vlakna, neprebojen material, nanodelci, degradacija
1 INTRODUCTION: THE WORLD OF
PROTECTIVE MATERIALS
Protection represents an important industry, both
economically and socially speaking, that includes, in the
broad sense, industrial, laboratory, home and, of course,
military protection, with an enormous variety of pro-
ducts, from simple plastic gloves to sophisticated and
confidential military armour. It is considered a "frag-
mented industry", in spite of the high volumes involved,
for literally thousands of manufacturers of raw materials,
producers of finished goods, distributors at all scales,
consultants, etc., operating worldwide with a steady
growth rate of around 3 % annually.1–10 In particular, the
materials that may be used for personal garment have
attracted a great deal of attention in the past few decades.
In fact, in December (2003), U.S. Attorney General John
Ashcroft instructed the National Institute of Justice to
implement a new initiative "to address the reliability of
body armour...and to examine the future of bullet-
resistant technology and testing…"2,11–17
In the particular area of protective materials, thanks
to the technological breakthrough of the 1960s that will
be discussed later on, the share of the market for high-
performance fibres for garment applications of one
single company (Dupont) is 60 %, whereas Honeywell
has 30 % and Toyobo, with the most recent technologies,
has 5 %.3–17
Nanotechnology is already a subject taught not only
at universities but also at the industrial level, while the
capacities of creating new nanomaterials have been
explored only in part.18–32
In the sections below we will focus, among all the
ballistic conceivable materials, on those intended for
personal protection, particularly on the ballistic vests and
related gadgets.
2 MODERN TRENDS IN BALLISTIC
MATERIALS
Today’s generation of body-armour systems can
provide protection at various levels designed to defeat
most common low- and medium-energy handgun rounds.
However, currently the highest-threat-level ballistic
needs in the market are fulfilled by special, high-
performance ceramics that tend to be very costly, fragile
for standard handling, extremely heavy and very difficult
to shape to the requirements of an ergonomic design.3–5
The search for novel polymer-based armour materials
dates back to the invention of synthetic macromolecules.
Accordingly, some companies have recently publicized
some nanofibre-reinforced systems that are expected to
provide very attractive weight/protection relationships
not only for personal equipment, but also for belly plates
for motor vehicles and even aircrafts5, exposed to the
impacts caused by dust, birds and other objects, not
necessarily by combat conditions, offering an interesting
potential market for novel ballistic materials.
From the military point of view and according to a
recent report3, nanotechnology offers two important
Materiali in tehnologije / Materials and technology 47 (2013) 3, 267–271 267
UDK 620.3:66.017 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 47(3)267(2013)
advantages: first, the potential to achieve high degrees of
miniaturization, which will be reflected in the weight of
the equipment and second, the possibility of finding
unexpected effects at the nanoscale, which not only will
represent a strategic advantage over the enemy, but will
also include a possibility of concealing the technology
behind a given effect.
According to this view, a list of potential applications
of nanotechnology for equipping the soldiers of the XXIst
century is limited only by imagination, at least according
to the Massachusetts Institute of Technology’s Institute
for Soldier Nanotechnologies, which includes 56 specific
projects divided into 7 work teams,3 one of them dedica-
ted to energy-absorbing materials, an area of obvious
relevance for the use of nanoparticles and nanostructured
composites, in addition to "smart" materials, of course4.
3 KEVLAR: A TRUE TECHNOLOGICAL
REVOLUTION
The modern history of ballistic polymers begins with
Stephanie Kwolek’s 1966 patent on Kevlar, a para-ara-
mid, invented while working with Dupont.6 Chemically
speaking, the "para" configuration allows the formation
of fibres, as opposed to the "cis" one that is sterically
hindered due to the large aromatic groups of the struc-
ture. Thus, the discovery of the properties of this mole-
cular configuration has led to the development of a
whole family of high-performance polymeric materials:
the so-called p-aramids. In particular, Kevlar certainly
represents a technological revolution not only for armour
materials, but for many other important applications,
from brake lining to space vehicles, including boats,
parachutes, building materials, etc. From the chemistry
standpoint, Kevlar is an aromatic polyamide, produced
with a condensation reaction of para-phenylenediamine
and terephthaloyl chloride, yielding a product with a
chemical composition of poly-para-phenylene terephtha-
lamide (PPD-T), having the technical name of Kevlar. It
is known that aromatic and amide groups of the type
contained in the structure of PPD-T provide a high
mechanical and thermal strength. One of the important,
and not very well known, facts about Kevlar is that it
constitutes a type of liquid crystalline polymer. Indeed,
when PPD-T solutions are extruded to produce an actual
fibre, the liquid crystalline nanodomains align them-
selves according to the flow, thus producing a highly
anisotropic material, capable of withstanding very high
impact energies. For example, the tensile modulus of
Kevlar 29, a high-toughness variant used for ballistic
vests, is of around 60 GPa, which can be further in-
creased to 130 GPa (Kevlar 49) with thermal treatments
under tension, increasing the anisotropy of the crystal-
lites in the material7. The aromatic rings in the structure
of Kevlar provide a high thermal stability, since the
corresponding decomposition temperature is nearly
430 °C.8 After the success of the original Kevlar formu-
lation, Dupont and a number of other companies have
developed a whole family of p-aramids that, along with
the other special polymeric materials (e.g., ultra-high-
molecular-weight polyethylene –UHMWPE), nowadays
constitute the core of the ballistic vest industry.
One very important limitation of Kevlar, however, is
its susceptibility to degradation due to UV exposure,
environmental humidity and the chemicals contained in
perspiration (sweat), the conditions that cannot be
avoided during in-field operations. A report by the U.S.
Lawrence Livermore Laboratory9 reveals that Kevlar "is
susceptible to photo-degradation from UV light sources".
Photo-degradation is a phenomenon, in which the tensile
strength of the fibres is reduced as a result of exposure to
UV light sources such as sunlight and fluorescent light.
Photo-degradation leads to reduced abrasion and tear
resistance in aramid fibres such as Kevlar. This problem
is so serious that, for example, the User Instruction,
Safety and Training Guide provided to the customers by
Lion Apparel (Dayton, Ohio, U. S.) gives the following
warning:
"Exposure to ultraviolet light (found in the sun’s rays
and fluorescent light) will severely weaken and damage
the fabrics in your protective clothing after only A FEW
DAYS. This is especially true for fabrics of the following
aramid materials: Hoechst Celanese Pbi, Dupont Kevlar,
Dupont Nomex, Dupont Nomex Omega, Dupont Nomex
IIIA, Lenzing P84, Southern Mills Advance, and BASF
Basofil."
4 NOVEL BALLISTIC FIBRES
One of the most promising recent advances in poly-
meric materials for protection garments is the Japanese
fibre Zylon10, a poly(p-phenylene benzobisoxazole)
(PBO), which has a tensile strength of around twice that
of Kevlar and similar commercial p-aramids, such as
Twaron (by Teijin Twaron Co.). The amazing properties
of Zylon have allowed super light, very comfortable (and
very expensive!) vests. However, recent studies9-12 reveal
inherent limitations in terms of its degradation under
visible light, heat and, particularly, when exposed to
humidity and the chemicals commonly found in sweat,
which can lead to a 65 % strength loss over a period of
only six months.10 Sealing the fabric into some thermo-
plastic does not improve much this effect, even at room
temperature, due to the capillarity behaviour of the Zylon
fibre.12 These important limitations, along with the price,
have severely limited the use of an otherwise attractive
material.
5 STRATUM nanoPROTECTM:
NANOTECHNOLOGY IN ACTION
In the above context, armour industry faces an
interesting conundrum: the availability of very strong
polymeric materials that are highly sensitive to standard
V. M. CASTAÑO, R. RODRÍGUEZ: NANOTECHNOLOGY FOR BALLISTIC MATERIALS ...
268 Materiali in tehnologije / Materials and technology 47 (2013) 3, 267–271
working conditions. A number of solutions have been
proposed and tested so far, from protective coatings, to
the use of irradiation13–16 to change the molecular confi-
guration of polymeric materials such as HIPS, Nylon,
Kevlar and Zylon and make them less susceptible to the
environment, to the humidity coming from both the
ambient and, particularly, from the bearer of the vest
(i.e., the sweat). The problem is not simple at all,
because the treatment or coating must not only preserve
the ballistic properties of the system, but also allow a
good deal of comfort that can be affected by stiffening
the fibres through crosslinking or surface layers of
various kinds. The challenge is, thus, too great for the
standard technology of ballistic materials, as it can be
corroborated by the fact that the leading industries in the
field have spent years and enormous amounts of money
to produce an environmentally stable garment to no
avail.
An opportunity for nanotechnology is then to find a
solution using the unique characteristics of nanoscale
systems. Indeed, ceramics are known to have a very high
resistance to UV degradation, as compared to poly-
mers.17–30 Their brittleness and specific density, however,
prevent their use in a garment. By chemically attaching
suitable ceramic nanoparticles to the surface of the fibre,
one is able to effectively shield the material against UV
without changing any other property. This can be done
with a proper reaction between the nanoparticles and the
previously modified Kevlar 29 fibre, using an organic
coupling agent for the fibre and the particles.31–34
The problem of humidity is a more complicated one,
since the common paradigm is to offer a physical barrier
to the water molecules. The obvious difficulty is to
ensure, at a molecular level, that the barrier keeps its
integrity during the use, enduring bending, shear and all
the typical abuses of a military garment. Chemically
functionalized nanoparticles linked to organic structures
offer a possibility of presenting a chemical barrier to
water molecules. This has a number of unique advan-
tages: first, no need to have a 100 % continuous coverage
of the surface; second, no danger of detachment as with a
coating, and third, no change in the other relevant
physical and chemical properties of the fibre.34
The above hypothetical scheme has been applied to
an actual armour vest, first commercially produced by
Parafly, S. A.,17 under the trade mark of STRATUM
nanoPROTECTM, and currently by QUANTICORP.33 The
proprietary technology includes a multicomponent
network, which uses, among other materials, Kevlar 29,
to which chemically modified nanoparticles were
attached during the fabrication process. Photographs of
the modified Kevlar 29 fabric can be appreciated in
Figures 1 and 2.
In the case of STRATUM nanoPROTECTM, the
Kevlar fibres were thermally treated to become ribbon-
like fibres that changed, after the weaving, into a cloth
with a better resistance to bullet penetration. This impro-
vement is essentially due to an increase in the resistance
to fibre openings in the fabric as a bullet penetrates the
vest. Figure 3 displays a photograph of a standard
V. M. CASTAÑO, R. RODRÍGUEZ: NANOTECHNOLOGY FOR BALLISTIC MATERIALS ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 267–271 269
Figure 3: Detail of a standard Kevlar fabric using threads with a
cylindrical shape
Slika 3: Detajl navadnega kevlarskega tkanja z uporabo niti z valjasto
obliko
Figure 1: Overview of Kevlar-based STRATUM nanoPROTECTM
fabric
Slika 1: Videz tkanine STRATUM nanoPROTECTM, ki temelji na
kevlarju
Figure 2: Detail of the fibres
Slika 2: Detajl vlaken
Kevlar fabric employing threads with a cylindrical shape,
where a clean hole produced by a projectile 9 mm is
clearly observed.
STRATUM nanoPROTECTM is then one of the first
commercial examples of an effective use of nanotech-
nology in armour devices. Nanotechnology added very
significant UV and chemical resistance to the ballistic
performance of the vest, which fulfils the international
regulations in the area, being a unique garment. Figure 4
shows a photograph of a STRATUM nanoPROTECTM
vest.
6 CONCLUDING REMARKS
Today, the area of materials for various types of
protection represents a unique opportunity for nanotech-
nology, though perhaps not to develop brand new
systems with amazing properties in a short term, but to
overcome some of the serious limitations of the current
technologies discussed above. Some other desirable
features for armour garments can be achieved with
current techniques available to many groups working in
nanosystems throughout the world. In particular, the
development of adequate variations of the nanotech-
nology products described above, specifically Zylon and
other high-performance ballistic fibres,35 is currently
under way and will be reported separately.
Acknowledgements
The authors are indebted to Dr. Domingo Rangel, Dr.
Genoveva Hernandez and Mrs. Carmen Vazquez for their
support in various stages of this project.
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V. M. CASTAÑO, R. RODRÍGUEZ: NANOTECHNOLOGY FOR BALLISTIC MATERIALS ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 267–271 271

V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE
ASSESSMENT OF DIFFERENT TOOL- AND
HIGH-SPEED-STEEL PROPERTIES
VE^FUNKCIJSKI KIc-PREIZKU[ANEC ZA DOLO^ANJE
RAZLI^NIH LASTNOSTI ORODNIH IN HITROREZNIH JEKEL
Vojteh Leskov{ek, Bojan Podgornik
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
vojteh.leskovsek@imt.si
Prejem rokopisa – received: 2013-01-31; sprejem za objavo – accepted for publication: 2013-04-02
Depending on the differences in the balanced alloy composition and steel processing technology, the properties of tool and
high-speed steel, like temper resistance, hot yield strength, ductility and toughness, thermal fatigue and shock resistance, as well
as wear resistance can differ considerably among the same type of steel. A high hot-yield strength, a high temper resistance and
a good ductility tend to result in a high resistance to thermal fatigue, while a resistance to mechanical and thermal shocks
depends on the ductility and toughness. However, the properties of tool and high-speed steels also depend on the final vacuum-
heat-treatment process. Normally, hardness and fracture toughness are used to determine the influence of vacuum-heat-
treatment parameters and to optimize them for the specific operating conditions of the tool. However, there are also other tool
properties which are equally important and need to be taken into consideration. To determine such a wide range of properties,
different test procedures and different test specimens are required since none of the standard tests alone is capable of providing
the relevant properties completely. Currently the best overall appraisal of tool and high-speed steel applicability seems to be a
combination of fracture toughness, bending or compression testing and in specific cases of impact or small-punch creep tests.
The aim of the paper is to show the possibility of using a single KIc-test specimen for the determination of such a wide range of
properties, being important for tool and high-speed steels. Besides that, the usability of KIc-test specimens for the assessment of
technological properties such as nitridability, machinability, wear resistance, etc. was confirmed.
Keywords: tool steel, vacuum heat treatment, characterization, fracture toughness
Odvisno od razlik v sestavi in tehnologije izdelave jekla lahko pri isti vrsti orodnega ali hitroreznega jekla mo~no variirajo
lastnosti, kot so: odpornost proti popu{~ni krhkosti, natezna trdnost v vro~em, duktilnost in `ilavost, odpornost proti termi~nemu
utrujanju in udarcem, kot tudi odpornost proti obrabi. Visoka meja elasti~nosti v vro~em, velika odpornost proti popu{~anju in
dobra duktilnost se izra`ajo v visoki odpornosti proti termi~nemu utrujanju, medtem ko sta odpornost proti mehanskim udarcem
in termi~nim {okom odvisni od duktilnosti in `ilavosti. Vendar pa so lastnosti orodnega in hitroreznega jekla odvisne tudi od
procesa kon~ne vakuumske toplotne obdelave. Navadno se uporabljata trdota in lomna `ilavost za ugotavljanje vpliva
parametrov vakuumske toplotne obdelave in za njeno optimiranje pri orodjih, namenjenih specifi~nim razmeram pri uporabi.
Vendar obstajajo tudi druge lastnosti orodja, ki so enako pomembne in jih je potrebno upo{tevati. Za dolo~anje tako {irokega
podro~ja lastnosti so potrebne razli~ne vrste preizkusov in razli~ni preizkusni vzorci, saj nobena standardna preizkusna metoda
ni sposobna ugotavljanja vseh lastnosti. Trenutno najbolj{e presojanje uporabnosti orodnih in hitroreznih jekel omogo~a
kombinacija lomne `ilavosti, upogibne ali tla~ne trdnosti in v posebnih primerih udarni preizkus ali "small-punch" preizkus
lezenja. Namen ~lanka je prikazati mo`nost uporabe KIc-preizku{anca za dolo~anje razli~nih lastnosti, pomembnih za orodna in
hitrorezna jekla. Poleg tega je bilo potrjeno, da KIc-preizkusni vzorec omogo~a tudi dolo~anje tehnolo{kih lastnosti, kot so
sposobnost za nitriranje, obdelovalnost, obrabna odpornost in druge.
Klju~ne besede: orodno jeklo, vakuumska topotna obdelava, karakterizacija, lomna `ilavost
1 INTRODUCTION
The forming industry is confronted with ever-increas-
ing demands for higher productivity, lower production
costs and more complex products, which together with
an increased focus on advanced, high-strength, low-
weight materials put increased requirements on tools and
dies.1–4 Consequently, this means strengthened property
requirements for the tool and high-speed steels, includ-
ing temper resistance, hot yield strength, ductility and
toughness, wear resistance, thermal fatigue and shock
resistance.5,6 Furthermore, with the tool design being
pushed to the very limit of the material strength, any
unexpected deterioration in the tool material properties
will eventually lead to premature and unwanted tool
failure. Therefore, tool and high-speed steels are conti-
nuously subjected to the development, both in the
direction of improved properties as well as better quality
with reduced properties’ deviation.
Many of the tool- and high-speed-steel grades used
today have been developed over a period of several
decades. The most significant developments to date are
the balanced chemical composition and the introduction
of specific, technologically optimised production steps
for an optimally annealed condition. Depending on the
differences in the balanced alloy composition and the
optimised processing route, the material properties can
differ considerably among the same type of martensitic
tool and high-speed steels. A high hot-yield strength, a
high temper resistance and a good ductility tend to result
in a high resistance to thermal fatigue. The resistance to
mechanical and thermal shocks depends mainly on the
Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283 273
UDK 539.42 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 47(3)273(2013)
ductility and toughness, but to some extent it is also
related to a high yield strength. Finally, the hardness and
microstructure will define the friction, the wear and the
anti-galling properties of the tool.
All the properties of tool and high-speed steels
depend not only on the balanced chemical composition
and the processing route, but also greatly on the final
vacuum-heat-treatment process, which defines the final
microstructure. Real metallic materials, i.e., tool steels,
usually have a multi-phase microstructure with carbide
precipitates and non-metallic inclusions. When the tools
are subjected to loads, local stress concentrations will
occur next to these hard, non-deformable, wear-resistant
particles, which, if stresses cannot be released through
micro-yielding of the matrix, will accelerate tool
breakage.7–9 Traditionally, a trade-off between a tough
matrix and hard, wear-resistant carbide precipitates was
required. Vacuum heat treatment, on the other hand,
allows the optimization of the tool-steel microstructure,
which satisfies ever greater demands on the properties of
tools and dies, particularly in respect of a greater fracture
toughness, while maintaining or even increasing the
hardness and wear resistance.10 In this respect the
hardness and fracture toughness, KIc, were found to be
the most suitable parameters when optimizing the final
vacuum heat treatment of tool and high-speed steels.11
Despite the enormous variety of tooling operations,
some basic properties of tool materials are common to
almost all applications. These properties are the ductility,
toughness and hardness.8,10,12 Ductility and toughness
prevent instantaneous fracture of the tool or tool edges
due to local overload, while a high hardness prevents any
local plastic deformation. However, ductility, toughness
and hardness are more or less mutually exclusive
properties, which means the prevention of instantaneous
tool failure is often connected with a critical hardness
level that must not be exceeded in a specific application.
On the other hand, a low hardness may lead to a prema-
ture thermal shock, despite the associated high toughness
level. Therefore, for a given hardness, the ductility and
toughness of the tool should be as high as possible to
ensure a good cracking resistance.10
Besides hardness, ductility and toughness there are
also other tool steel properties that are becoming equally
important as we move towards more and more complex
tools. These include creep and wear resistance, bending
strength, elastic properties under bending as well as
compressive properties at room and elevated tempera-
tures. Finally, the time and costs required to produce a
tool depend on the technological properties, like high-
speed machinability, grindability, nitridability, etc.
Although all these properties can be determined using
standard test methods, each one requires specific and
often quite unique test specimens, which are expensive
and not always easy to produce. Furthermore, different
geometries of standard specimens mean different heat-
treatment conditions, which makes it practically impossi-
ble to directly correlate the properties of tool and
high-speed steels after the final heat treatment.
Therefore, the aim of this research work was to
determine the applicability of a single KIc-test specimen
for the determination of a very wide range of properties
in the final heat-treated condition, which are important
for tool and high-speed steels. Furthermore, work is
intended to show the potential of KIc-test specimens for
optimizing the final vacuum heat treatment of tool and
high-speed steels.
2 RELEVANT TOOL STEEL PROPERTIES
2.1 Ductility, toughness and hardness
Ductility and toughness were found to be the most
relevant properties in terms of the resistance to total
failure of tool and high-speed steels, being a result of
mechanical or thermal overloading.8,12 However, these
are two different material properties, even though both
are too often referred to as toughness since the opposite
of both is brittleness. Unfortunately, no standardised
tests for the determination of toughness or ductility are
in common use and although some data are available, the
use of different test methods leads to confusion. The
importance of ductility and toughness for tool-steel
performance depends a lot on the tool geometry.12 In the
case of smooth un-notched surfaces the ductility and the
fracture stress are the relevant material properties.
However, if sharp notches or cracks are present, which is
more critical, the toughness is the most relevant property.
From this point of view, tool and high-speed steels
should be optimised in terms of ductility for the
un-notched regions, and in terms of toughness for the
notched regions.
The most reliable measure of toughness is the plain-
strain fracture toughness. The same value of fracture
toughness should be found for specimens of the same
material but with different geometries and with a critical
combination of crack size and shape and fracture stress.
Within certain limits, this is indeed the case, and infor-
mation about the fracture toughness can be used to
predict the failure for different combinations of stress
and crack size and for different geometries.13 Although
standardized methods for a fracture-toughness determi-
nation are available,14,15 their applicability is limited in
the case of hard and brittle materials, such as tool and
high-speed steels.10,11,16 Due to the high notch sensitivity,
the manufacture of a sharp fatigue crack is difficult and
expensive. On the other hand, a method based on a
circumferentially notched and fatigue-precracked tensile
specimen17,18 (Figure 1) has been found to be the most
promising alternative method.11,17,19 The fatigue crack in
the specimen can be obtained without producing any
disturbing effect on the fracture toughness of the steel if
such a crack is obtained in the soft-annealed condition,
i.e., prior to the final heat treatment.19
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
274 Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283
The advantage of the KIc-test specimen over the
standardized CT specimen (ASTM E399-90) lies in the
radial symmetry, which makes the specimen particularly
suitable for studying the influence of the microstructure
on the fracture toughness. The advantage relates to the
heat transfer, which provides a completely uniform
microstructure.19,20 Furthermore, in the case of KIc-test
specimens the fatigue crack can be created with rota-
ting-bending loading before the final heat treatment.11,19
And finally, for KIc-test specimens, plain-strain con-
ditions can be achieved using specimens with smaller
dimensions than those of conventional CT test speci-
mens. Shen Wei et al.21 proved that for circumferentially
notched and precracked round-bar tensile specimens,
plain-strain conditions are achieved when D  1.5 (KIc
/Rp0.2)2 and L  4  D (Rp0.2 is the yield stress).
In the case of circumferentially notched and fatigue-
precracked tensile-test specimens showing linearly
elastic behaviour up to fracture, the fracture toughness
KIc can be calculated using equation (1) 22:
K
P
D
D
dIc
= − +⎛⎝
⎜ ⎞
⎠
⎟
3
127 172. . (1)
where P is the load at failure, D is the outside non-
notched diameter, and d is the diameter of the instantly
fractured area, i.e., the diameter of the ligament next to
the crack. Eqn. (1) is valid as long as the condition 0.5 <
d/D < 0.8 is fulfilled. It is known that the fracture
toughness of conventionally produced tool steels
depends on the specimen orientation.23 It is the highest
for crack propagation perpendicular to the rolling
direction and the lowest for a crack propagating along
the rolling direction. Therefore, the way how specimens
are taken from the master block is very important and to
be on the conservative side specimens should be taken
in a short transverse direction.
Besides toughness also the hardness of tool and
high-speed steels greatly depends on the vacuum-heat-
treatment procedure, with both properties being mutually
related. The hardening mechanism is different for
as-quenched and fully-heat-treated tool steels. In the
as-quenched tool steel, mostly the work-hardening and
solid-solution hardening affect the steel’s hardness.
Tempering then leads to the precipitation of carbide
particles and a significant decrease of the carbon content
in a solid solution of martensite as well as of the density
of dislocations. Thus, the hardness of fully-heat-treated
tool steels is mainly the result of precipitation hardening
and, to a small extent, solid-solution hardening. A high
hardness is required to facilitate a high resistance to
plastic deformation and wear. However, with a high
hardness the toughness is reduced, which can lead to
cracking. Therefore, for a given hardness level, the
fracture toughness of tool and high-speed steels should
be as high as possible to ensure a good cracking resi-
stance.
2.2 Bending strength
The standard bend testing, which is widely used to
measure the bending fracture strength of tool and
high-speed steels is considered to be the most reliable
and to give a great deal of information regarding the
toughness and the ductility of the material. According to
the ASTM E290-09 standard,24 bend testing can be
performed under 3- or 4-point loading. The bending
moment diagrams show that for each mode compression
stress is present on the concave side of the specimen,
which for the 4-point mode is constant in a central span
between the two inner supports, the opposite to the peak
concentration under the central support of the 3-point
method. Stresses under bending can be calculated using
simple conventional theory, but only if the stress
throughout the specimen remains below the limit of
proportionality. If the limit of proportionality is exceeded
than the calculated stress will be higher than the actual
stress, which in the extreme case of fully plastic
deformation reaches 1.5 times the actual stress value.
Hoyle and Ineson25 showed that by progressively
correcting the calculated stress, taking into account the
limit of proportionality, a curve similar to the tensile
curves can be obtained. An analysis of such a corrected
curve can provide data related to the toughness
properties, and more important, the behaviour of tool and
high-speed steels. By performing experiments in the
elastic region of high-speed steel, it was found that the
test can discriminate between the elastic behaviour of
good and bad specimens.25 This can be quantified by
using an empirical relationship L/, the ratio of the
actual limit of proportionality L to the fracture stress ,
taken at a given hardness. Another parameter obtained
from the bending-test curve is the amount of plastic
deformation, or more correctly the deformation beyond
the limit of proportionality. This parameter is used to
determine the energy to fracture. Finally, with bend
testing the modulus of elasticity is also obtained.
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283 275
Figure 1: Circumferentially notched and fatigue-precracked KIc-test
specimen (all dimensions are in milimeters)
Slika 1: KIc-preizku{anec z obodno zarezo in utrujenostno inicialno
razpoko (vse dimenzije so v milimetrih)
2.3 Compression properties
For tool and high-speed steels in the hardened and
tempered condition the compression properties are
considered to be of far greater importance than the ten-
sile properties. Namely, their ductility in the hardened
and tempered condition is low and therefore not easily
revealed in a tensile test. The compression properties of
tool and high-speed steels at room and elevated tempe-
ratures can be determined using standard test methods
ASTM E9-09 26 and ASTM E209,27 respectively. The
parameters obtained from the compression test include
the compressive strength, the compressive yield strength
and the strain-hardening exponent. Compressive strength
correlates with the hardness of tool steel, while the
strain-hardening exponent describes its ductility.
2.4 High-speed machinability
The trend among toolmakers is to abandon the old
strategy of roughing and machining a piece of steel,
hardening it, finishing and grinding it, eroding features
impossible to cut, and polishing the surfaces. With the
introduction of high spindle-speed machines, which can
create an excellent surface finish that needs little or no
polishing, toolmakers prefer high-speed machining
(HSM) in hardened conditions. In the case of HSM com-
plex tool geometries can be machined in a single ope-
ration, thus eliminating many time-consuming operations
and reducing costs. Since HSM combines the roughing,
semi-finishing, and finishing operations of hardened
steel, spindles with high speed to allow fine step-overs
for finishing and with an adequate amount of torque for
roughing are required. Today, the HSM of hardened and
tempered tool and high-speed steels with an unin-
terrupted cut is possible up to a hardness of about 65–67
HRc.28
HSM is different from conventional machining and to
make it work a completely new philosophy needs to be
adopted.29 The understanding begins with the definition
of high-speed machining. Rather than using the conven-
tional definitions based on spindle speed, tool-tip speed,
or bearing ratio, it has to be defined in terms of fre-
quency. According to this, HSM occurs as the tooth-pass
frequency approaches a substantial fraction of the domi-
nant natural frequency of the machine-tool system.29 A
definition based on tool-pass frequency gives engineers a
tool for eliminating chatter and its deteriorating effect on
tool wear, surface finish, and machine life. Machinability
as well as the HSM of different tool and high-speed
steels can vary considerably. Factors largely influencing
the machinability of tool steels are the chemical compo-
sition, the microstructure, the hardness, non-metallic
inclusions and residual stresses. However, the main
factor generally regarded as influencing the machina-
bility of a tool steel is its hardness. The harder the
material is and the more carbides it has the more difficult
it is to machine.
2.5 Grindability
As grinding is an important operation in the final pro-
duction of many components from tool and high-speed
steel, the grindability of the material in the hardened and
tempered condition can have important consequences for
productivity and economics. Tool and high-speed steels
have a high hardness and wear resistance, which makes
them difficult to grind, especially when using alumi-
nium-oxide wheels. Furthermore, high temperatures
generated during the grinding of tool and high-speed
steels lead to thermal damage in the form of surface
oxidation, softening, tensile residual stress generation
and re-hardening.30 Therefore, one needs to differentiate
between thermal sensitivity, a material’s tendency to be
damaged at high temperatures, and grindability, the
material’s effect on the cutting ability of the abrasive
grits.31
Hardened and tempered tool and high-speed steels
are composed of tungsten-molybdenum carbides (1400
HV) and vanadium carbides (3300 HV) in a matrix of
martensitic steel. It was found that the vanadium content
and the size of the vanadium carbides are the dominant
factors affecting the grindability of high-speed steel.32
According to König and Messer33 grindability can be
expressed by the G-factor, i.e., the ratio of ground mate-
rial volume to grinding wheel volume lost. By taking the
equivalent carbide content (ECC), based on the stoichio-
metric sum of the carbide-forming elements in mass
fractions w (ECC/% = w(W) + 1.9 · w(Mo) + 6.3 · w(V)),
it was shown that the G-ratio is decreasing with
increased ECC.33 When comparing conventional and
powder metallurgy (P/M) tool steels the G-ratio for latter
can be up to 7 times higher, thus indicating better
grindability.34 In this case better grindability is the result
of a smaller carbide size found in the P/M steels.32
Besides the G-ratio, being an adequate way to mea-
sure grindability, the rate of power increase represents a
more meaningful measurement. Quantitatively, the
machinability and grindability can be determined on the
basis of a tribological evaluation as well as by measuring
the forces during cutting or grinding.35 Thus, the instru-
mented wedge grinding test (WGT),36 involving the
measurement of grinding forces in the tangential and
radial directions as well as the contact zone temperature,
gives the capability of characterizing the material
behaviour resulting from the grinding operation.
3 EXPERIMENTAL
3.1 Material and vacuum heat treatment
To demonstrate the applicability and potential of
KIc-test specimens (Figure 1), ESR high-speed steel
AISI M2 (delivered in the shape of rolled, soft-annealed
and peeled bars 20 mm × 4000 mm) with the following
chemical composition (mass fractions): 0.89 % C, 0.20 %
Si, 0.26 % Mn, 0.027 % P, 0.001 % S, 3.91 % Cr, 4.74 %
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
276 Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283
Mo, 1.74 % V, and 6.10 % W was used. KIc-test speci-
mens were cut from the delivered bars in the direction of
rolling and due to the high notch sensitivity of the
hardened high-speed steel, a fatigue pre-crack of about
0.3 mm was created prior to the final heat treatment
using a rotating-bending loading. All the KIc-test speci-
mens were then quenched in a horizontal vacuum
furnace, using N2 at a pressure of 5 bar. After the last
preheat the specimens were rapidly heated (25 °C/min)
to the austenitizing temperature of 1180 °C or 1230 °C,
soaked for 2 min and then gas quenched to a temperature
of 80 °C. Finally, high-pressure, gas-quenched KIc-test
specimens were double tempered for 1 h, which was
carried out in the same furnace at 9 different tempering
temperatures (Table 1). For each tempering temperature
at least 16 KIc-test specimens were prepared.
3.2 Material properties evaluation
The fracture-toughness measurement on the KIc-test
specimens was performed at room temperature using an
Instron 1255 tensile-test machine and a special specimen
fixture, which provided complete axiallity of the tensile
load. The cross-head speed was 1.0 mm/min, the speed
used for standard tensile-test specimens. During each
tensile test the tensile-load/displacement relationship
until failure was recorded, which for all KIc-test speci-
mens investigated showed linearly elastic behaviour, thus
confirming the validity of Eqn. (1).
After the tensile test the notch-section diameter d and
the radial distance of the crack initiation site from the
fatigue crack frontline x were measured for each frac-
tured surface using an optical microscope. The analysis
of the fractured surfaces was followed by a Rockwell-C
hardness (HRc) measurement, performed on each half of
the individual KIc-test specimen using a Wilson 4JR
hardness machine.
One half of the fractured KIc-test specimen was then
used to make a cylindrical 4-point bending-test specimen
(5 mm × 60 mm), with the bending test carried out
according to the ASTM E290-09 standard.24 The 4-point
bending-test specimens were prepared by high-speed
turning, using Sandvik SNMG 120408 K15 cutting
inserts, a feed rate of 0.1 mm/r, a depth of cut of 0.15
mm and a cutting speed of 100 m/min.28 After high-
speed turning, cylindrical specimens were further ground
by high-speed, centreless grinding (HSCG, vg = 63 m/s)
in order to obtain the prescribed average surface rough-
ness of 0.2 μm. During the preparation of the 4-point
bending-test specimens, high-speed turning and grinding
procedures were also used to evaluate the high-speed
machinability and grindability of the investigated high-
speed steel. The high-speed machinability was analysed
in terms of the evolution of the cutting inserts’ flank and
the notch wear and appearance of the chips. On the other
hand, the effectiveness of the grinding and grindability
were assessed on the basis of the final surface roughness,
achieved on 4-point bending-test specimens. In spite of
the fact that the G-ratio is normally used to measure the
effectiveness of the grinding wheel, we rather focused on
the quality of the ground surface.
The other half of the fractured KIc-test specimen was
cut 6 mm below the fractured surface in order to prepare
a metallographic sample for the fracture and microstruc-
ture examination. A further 12.5 mm down compression
test specimen (10 mm × 12.5 mm) was cut from the
fractured KIc-test specimen and tested on an Instron 1255
test machine at room temperature according to the
ASTM E9-09 standard.26 The compression-test results
included the compressive yield strength, the compressive
strength, the modulus of elasticity and the strain-harden-
ing exponent.
4 RESULTS
4.1 Fractography and metallography
Through an analysis of the fractured surface the main
crack-nucleation site can be identified. As shown in
Figure 2a, the main crack-nucleation site does not
coincide with the fatigue crack frontline in the case of
high-speed steels.11 For all the investigated specimens
the main crack-nucleation site was found to be slightly
away from the fatigue crack, with typical Chevron lines
designating the crack initiation and the direction of
propagation. Furthermore, no crack re-initiation in the
very tip of the existing fatigue crack could be observed.
A higher magnification of the crack-nucleation site,
denominated also as the weak spot, revealed the site to
be composed of large carbide clusters and strings and a
region of dimpled ductile fracture (Figure 2b).
The presence of the weak spot leads to an over
estimation of the fracture toughness. Blunting of the
fatigue-crack tip due to its branching out causes a
constraint effect and displaces the region of maximum
stresses and crack re-initiation from the area close to the
surface more into the bulk, ahead of the fatigue-crack
frontline.19 Consequently, the apparent fracture tough-
ness is elevated. In such cases a statistical analysis of the
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283 277
Table 1: Tempering temperatures used in the vacuum furnace
Tabela 1: Temperature popu{~anja v vakuumski pe~i
Tempering temperature (°C) 500 520 540 550 555 560 580 600 630
KIc-test specimens group A
TA = 1180 °C
A1 A2 A3 A4 – – A7 A8 A9
KIc-test specimens group B
TA = 1230 °C
B1 B2 B3 B4 B5 B6 B7 B8 B9
measured fracture toughness vs. the position of the weak
spot should be employed. If a linear correlation coeffi-
cient R is larger than 0.6 the fracture toughness obtained
on a statistically relevant number of specimens (>10) is
already the true value. Otherwise, a linear extrapolation
to the fatigue-crack frontline (i.e., x = 0) should be per-
formed.19 In the present case, for all specimen groups,
the linear correlation coefficient R was between 0.8 and
0.95, thus confirming the validity of the measured
fracture-toughness values.
Polishing the fractured surface and the preparation of
metallographic specimens reveal the microstructure of
the high-speed steel after a particular vacuum heat treat-
ment. Microstructures for KIC-test specimens austenitized
from 1230 °C and tempered at temperatures between
500 °C and 600 °C are shown in Figure 3. The micro-
structure of the investigated high-speed steel consists of
tempered martensite, fairly well dispersed undissolved
eutectic carbides and some stabilised retained austenite
in the matrix. The amount of retained austenite is
reduced with increased tempering temperature and
becomes nearly absent above 580 °C. A lower austeni-
zing temperature of 1180 °C was found to produce a
similar microstructure. However, it resulted in a higher
volume fraction of undissolved carbide particles, which
can be up to three times higher. Quantitative metallo-
graphy and microscopy, including Scanning Electron
Microscopy with Electron Back-Scattered Diffraction
(SEM/EBSD), further enable a detailed determination of
the eutectic carbide type, size and volume fraction, the
cumulative fraction of undissolved eutectic carbides and
carbide clusters, and the volume fraction of retained
austenite.
4.2 Fracture toughness and hardness
The fracture toughness and hardness values for two
austenitizing temperatures are presented in the tempering
diagram, shown in Figure 4. For both austenitizing
temperatures the fracture toughness shows a peak value
of more than 15 MPa m1/2 at the lowest tempering tem-
perature of 500 °C, which coincides with the relatively
high volume fraction of stabilised retained austenite and
the low hardness (60–61 HRc). By increasing the tem-
pering temperature the hardness level of the investigated
high-speed steel was increasing before reaching a peak
value of 64–66 HRc between 560 °C and 570 °C, which
was followed by a further decrease in the hardness.
However, at about the same hardness, the under-tem-
pered specimens quenched from the same austenitizing
temperature, show a higher fracture-toughness value. For
example, after vacuum quenching from 1230 °C and
double tempering at about 620 °C the investigated high-
speed steel achieves a hardness of 63 HRc and a fracture
toughness of 8.5 MPa m1/2. The same hardness, but with
an approximately 30 % higher fracture toughness, can be
obtained by using a tempering temperature of 510 °C
(Figure 4). The use of a lower austenitizing temperature
results in a reduced hardness and an increased fracture
toughness, with the difference becoming more evident at
higher tempering temperatures. These results clearly
show that the high volume fraction of stabilized retained
austenite in under-tempered high-speed steel signifi-
cantly improves its fracture toughness.
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
278 Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283
Figure 2: a) Typical fractured surface of a KIc-test specimen from
high-speed steel and b) main crack-nucleation site
Slika 2: a) Zna~ilna povr{ina preloma KIc-preizku{anca iz hitrorez-
nega jekla in b) mesto nukleacije razpoke
Figure 4: Tempering diagram for the investigated AISI M2 high-
speed steel
Slika 4: Diagram popu{~anja pri preiskovanem hitroreznem jeklu
AISI M2
Figure 3: Typical microstructure of group B KIc-test specimens auste-
nitized from 1230 °C
Slika 3: Zna~ilna mikrostruktura skupine B KIc preizku{ancev, avste-
nitiziranih iz 1230 °C
Based on the fracture toughness KIc and the ultimate
tensile stress u, estimated from the Rockwell-C hard-
ness, the fracture stress f can be calculated.37 A
reduction in the fracture stress is observed, when the
defect size, represented by carbides and carbide clusters,
exceeds a critical value.10 This gives us a tool to estimate
the critical defect size in vacuum heat-treated high-speed
steel depending on the austenitizing and tempering tem-
peratures. For example, in the case of an austenitizing
temperature of 1230 °C and a tempering temperature of
500 °C the critical defect size is in the range of 3–5 μm,
while it is reduced to about 1 μm for a tempering tempe-
rature of 560 °C, at the same time showing a faster drop
in the fracture stress with the defect size (Figure 5).
4.3 Bending and compression strength
The results of the bending and compression tests,
performed on 6 groups of specimens (A1, A4, A9, B1,
B4 and B9) are shown in Figures 6 and 7. The bending
strength of the investigated high-speed steel is increasing
for both austenitizing temperatures as the tempering
temperature gets higher. In the case of an austenitizing
temperature of 1180 °C the bending yield strength yB
and the bending strength B reached maximum values of
4000 MPa and 4150 MPa at the highest tempering
temperature of 630°C. By increasing the austenitizing
temperature to 1230 °C the bending yield strength and
the bending strength are reduced by about 15 %, as
shown in Figure 6.
The compressive strength of the vacuum heat-treated
high-speed steel shows the same dependency on the
tempering temperature as displayed by the Rockwell-C
hardness, with the peak values (yC 3000 MPa, C
3600 MPa) reached at the intermediate tempering
temperature of 550 °C. However, no noticeable diffe-
rence in the compressive strength was observed between
the two austenitizing temperatures used (Figure 7). On
the other hand, the strain-hardening exponent, indicating
the ductility of the material, shows a steep decrease with
increasing tempering temperature, while an increase in
the austenitizing temperature has the opposite effect, as
shown in Figure 7. The strain hardening exponent of the
investigated high-speed steel, austenitized from 1180 °C,
was reduced from 0.18 down to 0.04 as the tempering
temperature increased from 500 °C to 630 °C. When
using a higher austenitizing temperature of 1230 °C the
strain-hardening exponent also increased, indicating
better ductility. For under-tempered specimens (Ttemp =
500 °C) it increased to 0.31 and for specimens tempered
at 630 °C to 0.07 (Figure 7).
4.4 High-speed machinability and grindability
In order to evaluate the high-speed machinability
behaviour of vacuum heat-treated high-speed steel
investigation of cutting inserts was carried out. In the
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283 279
Figure 6: Bending-test results for vacuum heat-treated AISI M2
high-speed steel
Slika 6: Rezultati upogibnega preizkusa za vakuumsko toplotno obde-
lano hitrorezno jeklo AISI M2Figure 5: Example of the calculated fracture stress
Slika 5: Primer izra~unane napetosti loma
Figure 7: Compression test results for vacuum heat-treated AISI M2
high-speed steel; curves for c are obtained with quadratic regression
Slika 7: Rezultati tla~nih preizkusov vakuumsko toplotno obdelanega
hitroreznega jekla AISI M2; krivulje c so dobljene s kvadratno
regresijo
absence of a systematic observation of the cutting inserts
as a function of time the observed difference in the
appearance of the cutting inserts at the end of the tool
life was used to draw qualitative information. In most
cases a stable evolution of the wear along both the rake
and flank surface was observed, being a result of one or a
combination of the following modes; namely, abrasive
wear along with progressive micro-chipping, notching,
localized chipping and final destruction of the edge. A
comparison of the cutting inserts’ worn edge used to
machine specimens tempered at 500 °C (A1) and 550 °C
(A4), together with the machined surface and the corres-
ponding chips is shown in Figure 8. In the case of a
worse machinability tendency for notch formation and
adhesion of the work material to the cutting insert was
evident. However, when the work material adhesion was
postponed, as was the case for the specimens tempered at
550 °C (Figure 8b), the tool life was sufficiently pro-
longed (10 %). Indirect evidence of better machinabi-
lity also comes from the appearance of the chips, which
for the higher tempering temperature become more con-
tinuous, longer and smoother (Figure 8b).
From the investigation performed it can be concluded
that the high-speed machining of vacuum heat-treated
high-speed steel poses no technical problem. However, it
was observed that even at a comparable hardness level,
the increase in the carbide content observed for the lower
austenitizing temperature contributed to a deterioration
in the machinability of high-speed steel. At this point it
should be pointed out that proper cutting inserts
(material and micro-geometry) and cutting conditions
have to be selected to achieve the optimal tool life.
In spite of the fact that normally the G-ratio of a
grinding wheel is used to measure the grinding
efficiency, our focus was on the quality of the ground
surface. Therefore, the surface roughness of the 4-point
bending-test specimens (A1, A4, A9, B1, B4 and B9)
was analysed after grinding. The results are summarized
in Table 2.
Table 2: Surface roughness after high-speed centreless grinding
Tabela 2: Hrapavost povr{ine po "centreless" bru{enju z veliko
hitrostjo
Rough-
ness
4-point bending test specimen
A1 A4 A9 B1 B4 B9
Ra/μm 0.21 0.18 0.19 0.21 0.20 0.18
Rmax/μm 1.88 1.62 1.82 1.85 1.82 1.81
For the selected high-speed centreless grinding para-
meters and specimens austenitized from 1180 °C (group
A) the smoothest surface with a minimum average
roughness of 0.18 μm was reached in the case of speci-
men A4, tempered at 550 °C. In the case of specimens
austenitized from 1230 °C the surface quality was found
to improve with an increased tempering temperature, as
shown in Table 2. In terms of the roughness level, no
noticeable difference was observed between the two
austenitizing temperatures used. However, SEM micro-
scopy of the surface after high-speed machining and cen-
treless grinding revealed that for a higher austenitizing
temperature (group B) the undissolved eutectic carbide
particles were pulled out, which was not the case for
group A specimens, austenitized from 1180 °C.
5 DISCUSSION
The properties required from the tool and high-speed
steels greatly depend on the application and the process
the tool will be used for. In some instances the wear
resistance is the main concern, requiring a high hardness
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
280 Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283
Figure 8: High-speed machinability test results for: a) specimen A1 (t = 8.5 mm) and b) specimen A4 (t = 9.2 min); 1) edge of the cutting insert
at the end of the test, 2) machined surfaces and 3) associated chips
Slika 8: Rezultati preizkusa obdelovalnosti z veliko hitrostjo za: a) vzorec A1 (t = 8,5 mm) in b) vzorec A4 (t = 9,2 min); 1) rob rezalne plo{~ice
na koncu preizkusa, 2) obdelana povr{ina in 3) pripadajo~i ostru`ki
at sufficient fracture toughness, while in others the
fatigue properties and the resistance to crack propagation
are more important, demanding a high fracture toughness
and bending strength. Thus, the combination of different
tool properties becomes important and should be opti-
mized for a specific tool through proper vacuum heat-
treatment. Besides tempering diagrams displaying the
fracture toughness and hardness as a function of the
austenitizing and tempering temperatures (Figure 4)
diagrams combining other tool and high-speed steel
properties can also be prepared in order to relate the pro-
perties of the material to its performance. For example,
when looking for toughness, a properties combination of
fracture toughness and bending or compression yield
strength should be used (Figures 9a and 9b), and for
ductility, a combination of fracture toughness and strain
hardening exponent should be used (Figure 9c). How-
ever, an accurate comparison and correlation of different
material properties, as exemplified in Figure 9d,
combining the fracture toughness, the strain hardening
exponent and the bending yield strength is only possible
when different specimens are vacuum heat-treated under
identical conditions, not only in terms of temperature
and time, but more importantly in terms of heat transfer
and microstructure uniformity.
As already mentioned, due to the axial symmetry and
uniform heat transfer the KIc-test specimens are parti-
cularly suitable for studying the influence of vacuum
heat-treatment parameters on the microstructure of
metallic materials and consequently on their properties.
Since different test specimens can be made from the
same KIc-test specimen, a proper correlation between the
different material properties can also be carried out and
statistically analysed if using relevant number of KIc-test
specimens.
After measuring the fracture toughness, using the
KIc-test specimen, two halves of the test sample are
obtained, as shown in Figure 10. Both parts can be used
for a Rockwell-C or Vickers hardness measurement.
Then, from one of the halves a metallographic specimen
(10 mm × 6 mm) is cut for analysis of the fractured
surface and the microstructure just below the fracture.
V. LESKOV[EK, B. PODGORNIK: MULTI-FUNCTIONAL KIc-TEST SPECIMEN FOR THE ASSESSMENT OF ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283 281
Figure 10: Versatility of KIc-test specimen
Slika 10: Vsestranskost KIc-preizku{anca
Figure 9: Examples of tempering diagrams combining different properties required from tool material (a–c); a) bending yield strength and
fracture toughness, b) compression yield strength and fracture toughness, c) fracture toughness and strain hardening exponent, and d) diagram
combining fracture toughness, strain hardening exponent and bending yield strength
Slika 9: Primeri popu{~nih diagramov, ki kombinirajo razli~ne lastnosti, zahtevane od materiala orodja (a–c); a) meja te~enja pri upogibu in
lomna `ilavost, b) meja te~enja pri tlaku in lomna `ilavost, c) lomna `ilavost in eksponent napetostnega utrjevanja in d) diagram, ki kombinira
lomno `ilavost, eksponent napetostnega utrjevanja in mejo te~enja pri upogibu
From the same part of the KIc-test specimen, also the
compression test specimen (10 mm × 12.5 mm) is cut
and used to determine the compressive strength, the
compressive yield strength and the strain-hardening
exponent, correlating with the ductility of the steels. The
remaining can be used for an evaluation of the creep
resistance using the small-punch method (8 mm × 0.5
mm), the thermal conductivity, the wear resistance or for
an assessment of technological properties such as
nitridability, suitability for hard coating deposition, etc.
The other half of the fractured KIc-test specimen is used
to manufacture a 4-point bending test specimen (5 mm
× 60 mm) for an assessment of the bending strength.
During the manufacturing the high-speed machinability
and grindability in the hardened and tempered conditions
can also be evaluated. If bending properties are not
required a modified specimen with a circumferential
notch (r = 10 mm, z = 1 mm) for instrumented impact
ZR testing25 can be prepared. An instrumented impact
test provides the initial and maximum fracture force, the
total fracture time, the time to maximum fracture force,
the force-time diagram and the work used. The main
advantage of such an impact test lies in the possibility to
obtain a single fracture without shattering. This allows
further examination of the fractured surface, including
X-ray diffraction.
6 CONCLUSIONS
In the present paper the multi-functional KIc-test
specimen is presented. The idea is to minimize the costs
for tool and high-speed steel characterization. Due to the
fact that none of the standard tests alone is capable of
describing all the relevant properties, more than one test
is required. Besides this, the test results need to be
related to the service behaviour of the tool and high-
speed steels.
With the use of a KIc-test specimen, it is possible to
simultaneously assess basic properties such as hardness,
fracture toughness, bend and compressive strength etc.,
which can be directly correlated to the vacuum heat-
treatment parameters and microstructure obtained. In this
way the heat treatment of the tool and high-speed steels
can be optimized for a specific application. The other
properties such as wear and creep resistance, machina-
bility, grindability, nitridability, hard-coatings adhesion
etc., can be determined and correlated to the microstruc-
ture as well as between each other.
As a result of the relatively simple and economical
manufacturing of the axially symmetric KIc-test speci-
men, a sufficient number of specimens (10–20) for stati-
stical analyses can be produced, allowing an accurate
and reliable evaluation of the tool and high-speed steel
properties.
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Materiali in tehnologije / Materials and technology 47 (2013) 3, 273–283 283

M. CONRADI: NANOSILICA-REINFORCED POLYMER COMPOSITES
NANOSILICA-REINFORCED POLYMER COMPOSITES
POLIMERNI KOMPOZITI OJA^ANI Z NANOSILIKO
Marjetka Conradi
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
marjetka.conradi@imt.si
Prejem rokopisa – received: 2013-02-28; sprejem za objavo – accepted for publication: 2013-03-07
In the fast growing field of nanotechnology, polymer nanocomposites have become a prominent area of current research and
development. Silica/polymer nanocomposites are dominating the polymer and composite literature as well as a variety of
applications, many industrial products and other significant areas of current and emerging interest. This review will give a
general overview of the leading and most commonly used techniques and strategies for the preparation of both silica fillers and
silica/polymer nanocomposites, followed by a discussion of the main characterization methods, mechanical testing, properties
and applications. Typical examples of different systems will be reported and referred to the corresponding references for more
detailed descriptions.
Keywords: nanosilica, polymers, composites
V sklopu hitro razvijajo~e se panoge, imenovane nanotehnologija, so dobili polimerni nanokompoziti pomembno vlogo na
podro~ju raziskav in razvoja. Polimerni silicijevi dioksidni nanokompoziti dominirajo tako v literaturi, namenjeni raziskavam
polimerov in kompozitov, kot tudi v {tevilnih aplikacijah ter industrijskih produktih. V tem preglednem ~lanku se bomo
osredinili na splo{ni pregled vodilnih in najpogosteje uporabljenih tehnik ter strategij za pripravo tako silicijevih vklju~kov kot
tudi polimernih silicijevih dioksidnih nanokompozitov. Sledila bo razprava o glavnih metodah karakterizacije silika/polimernih
nanokompozitov, mehanskih preizkusov, o lastnostih ter njihovih aplikacijah. Predstavljeni bodo primeri razli~nih polimernih
silicijevih dioksidnih kompozitnih sistemov skupaj z ustreznimi referencami za podrobnej{i vpogled.
Klju~ne besede: nano silicijev dioksid, polimeri, kompoziti
1 INTRODUCTION
Polymer (nano)composites have been extensively
studied over a long period of time. They are generally
organic polymer composites mostly filled with inorganic
fillers. Their properties combine the advantages of the
inorganic filler material (i. e., rigidity, thermal stability)
and of the organic polymer (i. e., flexibility, ductility,
processability). However, the main advantage of these
composites is characterized by the volume fraction and
size of the fillers. If the fillers decrease in size from the
micro- to the nanoscopic scale, unique properties of
polymer nanocomposites are emphasized as the small
size of the fillers leads to a dramatic increase in the
interfacial area as compared with the ordinary com-
posites. This interfacial area then creates a significant
volume fraction of the interfacial polymer with the
properties different from the bulk polymer even at low
filler loadings.1,2
Reinforcement agents such as glass particles,3,4
ceramic particles,5 layered silicates,6–8 metal particles,9
rubber plastics10 and thermoplastics11,12 have already
been successfully used. Several researchers have studied
the effects of the particle size and volume fraction on the
mechanical response of polymer composites.13–17 There
are a few analytical models based on the crack propa-
gation along the particle surfaces taking the interspacing
between the particles into account18 as well as mathe-
matically considered trapping, pinning and bridging of
the crack front on the particles.19–21 It has been shown
that the fracture phenomena of the composites filled with
nanometer-sized particles differ from the behaviour of
the composites filled with micrometer-sized or larger
filler particles.
However, among the numerous polymer composites,
silica/polymer nanocomposites are the most commonly
reported in the literature and are also employed in a
variety of applications, such as electronics, automotive
and aerospace industries as well as used in many
industrial products due to their good mechanical cha-
racteristics22. In order to further improve the properties
of silica/polymer nanocomposites, their preparation,
characterization, mechanical properties and applications
have become quickly expanding fields of research in the
past few years.
The aim of this review is to give a general overview
of the leading and most commonly used techniques and
strategies for the preparation of both silica fillers and
nanosilica/polymer composites, followed by a short dis-
cussion of the main characterization methods, mecha-
nical testing, properties and applications. Typical
examples of different systems are reported and referred
to the corresponding references for more detailed
descriptions.
2 PREPARATION OF SILICA FILLERS
Silica particles normally exist in a form of a fine,
white amorphous powder or colloid suspension. Its most
Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293 285
UDK 678.84:66.017 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 47(3)285(2013)
important characteristic is an extremely large surface
area and a smooth nonporous surface, which can
promote a strong physical contact when embedded in a
polymer matrix.
Nowadays silica particles are commercially available
in all sizes ranging from nanometer to micrometer;
however, several researchers still synthetize particles on
their own. Two main methods have been developed for
silica-particle formation: the sol-gel method and the
microemulsion method.23 In 1968, however, Stöber and
Fink24 introduced a simple synthesis of monodisperse
spherical silica particles starting with tetraethyl ortho-
silicate (TEOS 98 %), deionized water, ammonia (25 %)
and absolute ethanol (99.9 %) as the alkoxide precursor,
hydrolyzing agent, catalyst and solvent. In the process,
two mother solutions are prepared, one containing
ammonia–water, and the other containing TEOS–etha-
nol. The two solutions are mixed in a thermostatically
controlled water bath (50 ± 1) °C. After 60 min, the
resulting spheres are separated from the liquid phase
with centrifugation and then ultrasonically dispersed in
deionized water. Finally, the particles can be dried in an
oven at 50 °C to obtain white powder. Note that using
this method, the final particle size critically depends on
the reagent concentrations, molar ratio and reaction
temperature. As shown in Figure 1 good monodispersity
of silica spheres can be obtained with this method.25
The dispersion of silica fillers and, consequently, the
compatibility between the polymer and silica have a
crucial impact on the mechanical properties of silica/
polymer composites. As most of the polymers are
hydrophobic in nature, it is important to improve the
interfacial interaction between the matrix and silica via
silica-surface modification, which can also improve its
dispersion in the matrix at the same time. In general, the
surface of silica fillers can be successfully modified with
either chemical or physical methods.
Modification of silica via a chemical interaction
involves a modification of its surface with modifying
agents (i. e., silanes) or grafting polymers. The most
common way of making silica hydrophobic and polymer
compatible, is silanization. We can find a long list of
silane coupling agents that generally have hydrolysable
and organofunctional ends and can be represented as
RSiX3. X stands for hydrolysable groups, typically
chloro, methoxy or etoxy gropus, and R stands for the
organo group that has to be chosen according to the
properties of the polymer. Silanes are attached to the
silica surface through the reaction of hydrolysable
groups with the hydroxyl groups on the silica surface
while the alkyl chains interact with the polymer26
(Figure 2).
Hydrophobicity, on the other hand, can be success-
fully increased by grafting the polymer chains to silica
particles either with a covalent attachment of end-func-
tionalized polymers to the surface or with an in-situ
monomer polymerization.
If silica is modified via a physical interaction, the
procedure usually involves surfactants or macromole-
cules adsorbed onto its surface. In principle, a polar
group of surfactants is adsorbed to the surface of silica
by an electrostatic interaction. As a consequence, the
physical attraction between the silica particles within
agglomerates is reduced, making silica particles easy to
incorporate into a polymer matrix.27
M. CONRADI: NANOSILICA-REINFORCED POLYMER COMPOSITES
286 Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293
Figure 2: Schematic presentation of the silica-surface modification by
trisilanol heptaisobutyl silesquioxane (IB5(SiO3/2)8(OH)3) via a cova-
lent bonding26
Slika 2: Shematski prikaz modifikacije silicijeve povr{ine s trisilanol
heptaisobutil sileskvioksanom (IB5(SiO3/2)8(OH)3) na osnovi kova-
lentne vezi26
Figure 1: Scanning electron microscopy (SEM) images of SiO2
particles prepared according to the Stöber procedure, a) 650 nm silica
particles, b) 240 nm silica particles25
Slika 1: SEM-slika SiO2 delcev, sintetiziranih po metodi Stöberja, a)
silicijevi delci 650 nm, b) silicijevi delci 240 nm25
3 COMPOSITE PREPARATION
The main concern in composite preparation is the
mixing process and the obtained homogeneous silica
dispersion in the polymer matrix. Therefore, prior to the
composite preparation, the compatibility between the
two components has to be assured to avoid silica
agglomeration, using one of the methods described in the
previous section.
The simplest method of silica/polymer composite
preparation is direct mixing of silica into the polymer
matrix, i. e., by melt or solution blending. Besides blend-
ing, sol-gel processes and in-situ polymerization are also
widely used among the preparation techniques.
3.1 Blending
Melt blending is the most commonly used method in
composite preparation due to its efficiency and opera-
bility. In the process, the polymer and the inorganic filler
(i. e., silica) are sheared in the melt at a temperature
equal or greater than the melting point of the polymer.
Under suitable conditions the material exfoliates and
disperses to the desired extent. This technology is very
versatile and can be applied to various polymers.28–31 It is
also possible to add swelling and compatibilizing agents
in order to improve the exfoliation and reach a better
adhesion between the two major components.
Solution blending, on the other hand, is a liquid-state
powder-processing method that allows a good molecular
level of mixing. Solution blending can be achieved by
either dissolving only the polymer matrix or dissolving
both the matrix and the nanoparticles.
3.2 Sol-gel process
The sol-gel process is a synthesis route consisting of
the preparation of a sol, the successive gelation and the
solvent removal. Within the past decades sol-gel
processes have been widely used to synthesize novel
organic/inorganic composite materials. In the case of
composites, the goal is to carry out the sol-gel reaction in
the presence of polymeric molecules (i.e., the organic
phase) containing functional groups that improve their
bonding to the inorganic phase. This is a very successful
reinforcement technique that can generate filler particles
within a polymer matrix.
3.3 In-situ polymerization
In-situ polymerization is a very effective and fast way
to construct a nanocomposite material. In this method,
the fillers are first pretreated with appropriate surface
modifiers and then added directly to the liquid monomer
during the polymerization stage. Using the solution
method, fillers are added to a polymer solution using
solvents such as toluene, chloroform and acetonitrile to
integrate the polymer and filler molecules.
4 CHARACTERIZATION
A homogeneous distribution of finely dispersed
fillers in a silica/polymer composite is a prerequisite for
obtaining good mechanical properties of the end
material. In a composite, interfacial interaction between
the fillers and the polymer matrix plays a crucial role in
toughening the composite. In order to achieve a high
particle/polymer area and to distribute the mechanical
stresses within the composite, the composite has to
consist of homogeneously distributed filler particles
which are not agglomerated. Therefore, prior to further
mechanical testing, the chemical structure, morphology
and microstructure of the composites must be analysed.
4.1 Infrared and Raman spectroscopy
Infrared and Raman spectroscopy are normally used
to confirm the surface modification of the silica fillers
implemented in a polymer matrix. Building up a modi-
fier layer around the silica particles can be followed by
inspecting the spectra of the filler32,33 or the vibrational
bands of the surface modifier, the intensities of which
depend on the number of the filler surface sites occupied
by the modifier.
In Figure 3, typical IR transmission spectra of the
silica-surface modifier trisilanol (IO7) POSS, polymer
polyvinyl chloride (PVC), 30 nm silica/PVC and 130 nm
silica/PVC composites are presented.34 The silanol
groups of open-cage POSS do not spontaneously lead to
condensation, as silanols of other simple alkoxysilanes,
but preferentially interact with the active sites on the
surface of the filler.35 Normally, the bands attributed to
trisilanol POSS are weak because the Si-O-Si bands are
covered by the same band of silica spheres. However, a
closer inspection of the peak intensity of the band at
2907 and 2953 cm–1 indicates a progressive occupation
of the accessible sites on the silica surface. A compa-
rison of the infrared spectra in the C-H region of
trisilanol POSS (Figure 3 zoom), modified silica/PVC
M. CONRADI: NANOSILICA-REINFORCED POLYMER COMPOSITES
Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293 287
Figure 3: IR transmission spectra of trisilanol (IO7) POSS, PVC, 30
nm silica/PVC and 130 nm silica/PVC composites34
Slika 3: IR-transmisijski spektri trisilanola (IO7) POSS, PVC-ja,
kompozitov silicijev dioksid-PVC 30 nm in silicijev dioksid-PVC 130
nm34
composite and bulk PVC revealed the presence of
trisilanol POSS on the silica surface. The C-H band at
2953 cm–1 was also present in the spectrum of the bulk
PVC and could be seen in the spectrum of the PVC
composite as well. Surprisingly, the C-H band at 2907
cm–1 surpasses the band at 2953 cm–1 because of the
presence of isooctyl groups on the surface.
In Figure 4 typical Raman spectra of composite
silica/PVC recorded as a function of the size of silica
spheres is presented. In the Raman spectrum of PVC,
polarized lines are observed at 639, 695, and 1435 cm–1.
The band at 1435 cm–1 is assigned to the CH2 scissors
vibration and a broad doublet-like band envelope in the
region of 600-700 cm–1 was attributed to C-Cl stretching
vibrations.36 For comparison, we recorded the Raman
spectrum of pure, bulk, 600 nm silica spheres, which
have a very weak response in the fingerprint range (i.e.,
500–2000 cm–1) and also in the CH region. Inside the
composites under investigation, we, therefore, used the
mapping of silica spheres. The average spectrum of silica
mapping in Figure 4 indicates a modification of the
silica spheres as the spectrum includes the bands of pure
silica and the bands of trisilanol IO7 POSS. In addition,
PVC has a similar response in the CH region. A
successful modification of 30 nm silica/PVC and 130 nm
silica/PVC composites was further identified with an
elevated band at 2912 cm–1. This band reveals the
isooctyl structure of POSS mostly because the other
bands are not so strong in the Raman spectra.
4.2 Scanning electron microscopy and transmission
electron microscopy
Scanning electron microscopy (SEM) and trans-
mission electron microscopy (TEM) are the most
common methods used in the morphology evaluation.
In Figure 5 we see typical SEM micrographs of
POSS-silanated, 130 nm silica/PVC composites. We can
see that a 130 nm silica/PVC composite consisted of
3D-ordered silica particles. SEM micrographs also
revealed (Figure 5b) that the 3D silica-particle structure
was ousted towards the film surface, frequently
encountered with polymers consisting of two phases with
a poor mutual compatibility due to their different
chemical compositions. Typical examples are polymers
with a low surface energy obtained with an addition of
fluoropolymers or fluorosilanes.37
Another example of SEM imaging in composite-
morphology characterization is an analysis of the
fracture surface of a composite filled with a very low
amount of silica fillers, i.e., less than volume fraction
1 %. Due to the small amount of silica fillers it is not
possible to observe isolated particles or an arrangement
of particles as shown previously in Figure 5; however,
the information on the particle inclusion can be revealed
with an analysis of the properties of fracture surfaces.
In Figure 6 typical fracture surfaces of the neat
epoxy and diglycidyl ether of bisphenol A surface-
treated, 30-nm and 130-nm silica/epoxy composites are
presented.38 Prior to the SEM imaging the samples were
frozen in liquid nitrogen and broken by hand in order to
observe the natural crack propagation in the composite.
The inclusion of silica fillers in the epoxy matrix is
confirmed by an increased roughness of the composite’s
M. CONRADI: NANOSILICA-REINFORCED POLYMER COMPOSITES
288 Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293
Figure 4: Raman spectra of pure 600 nm silica spheres, trisilanol
(IO7) POSS, the average spectrum of silica mapping, pure PVC, 130
nm silica/PVC and 30 nm silica/PVC composites34
Slika 4: Ramanski spektri neobdelanih 600 nm silicijevih delcev,
trisilanola (IO7) POSS, povpre~nega spektra silicijevega mapiranja
~istega PVCja ter 130 nm silika/PVC in 30 nm silika/PVC kompo-
zitov34
Figure 5: a) Top-view and b) side-view SEM micrographs of a
silica/PVC composite composed of volume fraction 30 % of POSS-
silanated, 130 nm silica fillers34
Slika 5: SEM-pogled: a) z vrha in b) pogled s strani kompozitov
silicijev dioksid-PVC, sestavljenih iz volumenskega dele`a 30 %
POSS-silaniziranih silicijevih dioksidnih vklju~kov 130 nm34
fracture surface as compared to the smooth surface of the
pure epoxy. Both silica and epoxy composites break in
sharp fracture lines and characteristic steps decorated
with a fish-skin-like microstructure (Figure 6d) that also
indicates an increased brittleness compared to the pure
epoxy. The roughness of the fracture surface, however,
slightly decreased with the decreasing particle size.
Transmission electron microscopy can serve as a very
useful tool for determining the exact size and distribution
of the embedded nanoparticles. In Figure 7 a typical
micrograph of the metal nanoparticles embedded in a
polymer matrix is presented.
5 MECHANICAL PROPERTIES
One of the primary reasons for adding inorganic
fillers to polymers is to improve their mechanical
performance through optimization of the balance
between the strength/stiffness and the toughness of a
composite. The mechanical response of composites
strongly depends also on the silica content and size and
is generally characterized with respect to different
properties, such as the tensile strength, flexural strength,
hardness, impact strength, fracture toughness, etc.
5.1 Tensile-strength test
With the tensile test we analyse stress  and strain 
that are determined from the measured load and deflec-
tion using the original specimen cross-sectional area S0
and length l0 as follows:
 = F/S0,  = dl/l0
Typical stress-strain curves of pure PVC and
POSS-silanated silica/PVC composites with two types of
silica fillers, 130 nm and 30 nm in diameter, are summa-
rized in Figure 8. For silica/PVC composites, the trend
reported in Figure 8 can be divided into four stages
before the material breaks: the initial linear elasticity,
nonlinear transition to global yield, region of necking
and strain softening. However, the stress-strain behaviour
of pure PVC is different, typically with a broad constant
stress regime for increasing the strain without a pro-
M. CONRADI: NANOSILICA-REINFORCED POLYMER COMPOSITES
Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293 289
Figure 6: a) Hand broken fracture surfaces of pure epoxy, b) composite with the volume fraction 0.5 % diglycidyl ether of bisphenol A
surface-treated, 30-nm silica fillers, c) composite with 0.5 % diglycidyl ether of bisphenol A surface-treated, 130-nm silica fillers and d) a
fracture surface detail – a fish-skin-like microstructure – of the 130-nm silica/epoxy composite38
Slika 6: a) Lomna povr{ina vzorcev ~istega epoxyja, b) kompozita, obogatenega z volumenskim dele`em 0,5 % z diglicidil etrom bisfenola A
povr{insko obdelanih 30-nm silicijevih vklju~kov, c) kompozita, obogatenega z volumenskim dele`em 0,5 % z diglicidil etrom bisfenola A
povr{insko obdelanih silicijevih dioksidnih vklju~kov 130 nm in d) detajl lomne povr{ine – mikrostruktura ribje ko`e v kompozitu silicijev
dioksid-epoksi 130 nm38
Figure 7: TEM micrograph of the Fe2O3 nanoparticles embedded in
the polypropylene (PP) matrix
Slika 7: TEM-slika Fe2O3 nanodelcev v mre`i polipropilena (PP)
nounced necking or strain softening before the sample
breaking.
In the low strain portion of the curves all the
investigated samples followed Hooke’s law giving us the
information on the composite modulus of elasticity E. As
shown in Table 1, we observed a 30–40 % increase in E,
making 30 nm and 130 nm silica/PVC composites stiffer.
This large increase in elastic modulus is due to a large
amount of silica particles. Another material proportion
that we can obtain from the stress-strain diagram is the
maximum tensile strength (UTS). A significant
strengthening is observed in both 30 nm and 130 nm
silica/PVC composites, 20–30 %, respectively (Table 1).
An opposite response is, however, observed for the
elongation at break depending on the silica size in the
PVC matrix (Table 1). In contrast to the increased
stiffness of 30 nm and 130 nm silica/PVC composites,
their elongation at break that is lower than for pure PVC
(15–30 %) indicates an embrittlement effect upon an
addition of silica fillers.
5.2 Three-point bending test
The three-point bending test (3PB) covers the
determination of flexural properties of a material by
measuring the deflection of a sample under applied load.
Figure 9 shows a typical stress-strain curve for the
samples under investigation, diglycidyl ether of bisphe-
nol A surface-treated silica composites and the neat
epoxy. Table 2 lists the corresponding material pro-
portions obtained with the 3PB test: elastic modulus (E),
maximum tensile strength (UTS) and elongation at break.
We observed an approximately 10–20 % increase in
E and UTS for both composites compared to the pure
epoxy. The experimental scattering of both the Young’s
modulus and the UTS was less than 10 %. The incorpo-
ration of silica fillers, on the other hand, caused a
decrease in the elongation at break, which implies an
increase in the composite brittleness.
5.3 Fracture toughness
The fracture-toughness test is designed to
characterize the toughness of a material in terms of the
critical-stress-intensity factor, KIC, and the energy per
unit area of a crack surface or the critical strain energy
M. CONRADI: NANOSILICA-REINFORCED POLYMER COMPOSITES
290 Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293
Figure 9: Typical stress-strain curves of diglycidyl ether of bisphenol
A surface-treated silica/epoxy composites and the neat epoxy as
obtained with the 3PB test38
Slika 9: Z upogibnim preizkusom (3PB) dobljen zna~ilen diagram
napetost – deformacija z diglicidil etrom bisfenola A silaniziranih
kompozitov silicijevega dioksida-epoksi in ~istega epoksija38
Table 1: Elastic modulus (E), tensile strength (UTS) and elongation at
break of POSS-silanated silica/PVC composites and pure PVC eva-
luated with the tensile test34
Tabela 1: Z nateznim preizkusom dobljen elasti~ni modul (E), maksi-
malna natezna trdnost (UTS) in raztezek pri pretrgu POSS-silani-
ziranih kompozitov silicijev dioksid-PVC in ~istega PVC34
Sample E/GPa UTS/MPa elongation atbreak (%)
PVC 1.7 89.9 7.0
PVC + 60 % vol. frac-
tions of SiO2 30 nm
2.4 112.1 6.0
PVC + 60 % vol. frac-
tions of SiO2 130 nm
3.0 122.2 4.9
Table 2: Elastic modulus (E), tensile strength (UTS) and elongation at
break of diglycidyl ether of bisphenol A surface-treated silica/epoxy
composites and the neat epoxy evaluated with the 3PB test38
Tabela 2: Z upogibnim preizkusom (3PB) dobljen elasti~ni modul (E),
maksimalna natezna trdnost (UTS) in raztezek pri pretrgu z diglicidil
etrom bisfenola A silaniziranih kompozitov silicijev dioksid-epoksi in
~istega epoksija38
Sample E/GPa UTS/MPa elongation atbreak (%)
Epoxy 2.6 127 10.0
Epoxy + 0,5 % vol. frac-
tions of SiO2 130 nm
3.0 141 9.6
Epoxy + 0,5 % vol. frac-
tions of SiO2 30 nm
2.8 138 9.0
Figure 8: Typical stress-strain diagram for pure PVC and silica/PVC
composites with POSS-silanated, 130 nm and 30 nm silica fillers as
obtained with the tensile-strength test34
Slika 8: Z nateznim preizkusom dobljen zna~ilen diagram napetost-
deformacija ~istega PVCja in kompozitov silicijev dioksid-PVC,
obogatenih s POSS-silaniziranimi 130 nm in 30 nm silicijevimi
dioksidnimi vklju~ki34
release rate, GIC, at the fracture initiation. In the experi-
ment, the load-deflection curves are measured as shown
in Figure 10. The fracture toughness, KIC, is then calcu-
lated by using linear elastic fracture mechanics:39
K
SP
BW
fIC
c
/= 3 2 ( )	
where
{ }
f ( )
. ( )( . .
( )( )
/ )
	
	 	 	 	 	
	 	
=
− − − +
+ −
3 199 1 215 3 93 2 7
2 1 2 1
1 2 2
3 2/
	 =
a
W
0
S and Pc are the span length and the maximum load; B,
W and a0 are the thickness, the width and the pre-crack
length of the specimen.
In Figure 10 diglycidyl ether of bisphenol A
surface-treated silica composites and the neat epoxy
show a linear response until the brittle fracture occurs.
The extracted experimental results for the fracture
toughness, KIC, for the composites and the neat epoxy are
listed in Table 3, showing a fracture toughness increase
by 25–30 % with the addition of silica fillers.
Table 3: Fracture toughness of diglycidyl ether of bisphenol A sur-
face-treated silica composites and the neat epoxy38
Tabela 3: Lomna `ilavost z diglicidil etrom bisfenola A silaniziranih
kompozitov silicijev dioksid-epoksi in ~istega epoksija38
Sample KIC/(MPa m1/2)
Epoxy 0.66 ± 0.05
Epoxy + 0.5 % vol. fractions of
SiO2 130 nm
0.91 ± 0.06
Epoxy + 0.5 % vol. fractions of
SiO2 30 nm
0.93 ± 0.06
5.4 Charpy impact strength test
The Charpy test is used to evaluate the amount of
absorbed energy by a material during a fracture, there-
fore giving us information on the impact toughness of
the material. An example of the Charpy impact strength
test results is in Table 4, presenting the impact energy
and the impact resistance for diglycidyl ether of
bisphenol A surface-treated silica composites and the
neat epoxy.
Table 4: Impact energy and impact resistance of diglycidyl ether of
bisphenol A surface-treated silica composites and the neat epoxy38
Tabela 4: Energija pri prelomu in odpornost proti prelomu z diglicidil
etrom bisfenola A silaniziranih kompozitov silicijevega dioksida-
epoksi in ~istega epoksija38
Sample Impact energyEimp/J
Impact resistance
Rimp/(kJ/m2)
Epoxy 0.19 ± 0.02 6.4 ± 0.7
Epoxy + 0.5 % vol. frac-
tions of SiO2 130 nm
0.26 ± 0.02 8.9 ± 0.6
Epoxy + 0.5 % vol. frac-
tions of SiO2 30 nm
0.33 ± 0.03 10.8 ± 0.7
The addition of silica particles increases the impact
resistance as well as the impact energy up to 60 %.
Surprisingly, the results of the Charpy impact test are
also strongly influenced by a particle diameter.
6 APPLICATIONS
As shown earlier, nanosilica-reinforced polymer
composites significantly improve the mechanical pro-
M. CONRADI: NANOSILICA-REINFORCED POLYMER COMPOSITES
Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293 291
Figure 10: Load-deflection curves of diglycidyl ether of bisphenol A
surface-treated silica composites and the neat epoxy38
Slika 10: Diagram sila – deformacija z diglicidil etrom bisfenola A
silaniziranih kompozitov silicijev dioksid-epoksi in ~istega epoksija38
Figure 11: SEM micrographs of POSS-silanated, a) 600 nm and b) 30
nm silica/PVC composite coatings on the AISI 316L surface45
Slika 11: SEM-slike POSS-silaniziranih kompozitnih prevlek silici-
jevega dioksida-PVC a) 600 nm in b) 30 nm na povr{ini AISI 316L
jekla45
perties of the end material and exhibit some unique
properties allowing many potential applications. Silica/
polymer nanocomposites have been reported to be used
in coatings,40,41 optical devices,42 electronics,43 photo-
resist materials,44 etc.
In most cases silica/polymer nanocomposites are
used as protective coatings, either to improve the
mechanical characteristics of the substrate material (i. e.,
wear, scratch, abrasion resistance) or to insure the
corrosion resistance in various environments. Corro-
sion-protection nanosilica/polymer coatings on metallic
substrates, for example, provide an effective physical
barrier between the metal and its environment containing
aggressive species, such as enhanced chloride-ion con-
centration, O2 or H+. A uniform dispersion of nano-
particles (i.e., SiO2) in a chosen polymer matrix with
desirable characteristics (i. e., epoxy) is shown to
increase the surface hydrophobicity (i.e., the self-clean-
ing effect) and to improve the adhesion between the
composite coating and the metallic surface.26,45–47
In Figures 11a and b we can see an example of a
POSS-silanated silica/PVC coating adsorbed on a steel
substrate of type AISI 316L. Figure 12 further presents
potentiodynamic measurements of these coatings in a
3.5 % NaCl solution at room temperature indicating an
improved anticorrosion behaviour in a chloride-ion-rich
environment as compared to the clean AISI 316L
surface. Potentiodynamic curves reflect the decreased
corrosion-current densities and corrosion potentials for
the silica/PVC-coated AISI 316L substrates.
7 SUMMARY AND OUTLOOK
A modification of a polymer matrix with silica fillers
allows significant increases in the modulus and strength
contributions of the matrix to the overall composite
properties. Obtaining the optimum properties for the
nanocomposites, however, requires an excellent homo-
geneous dispersion of the fillers, as the tendency of the
silica particles to agglomerate can seriously affect the
achievable properties. Therefore, to provide a strong
interfacial interaction between the inorganic particles
and the polymer matrix, silica fillers must have suitably
modified surfaces. The end result is a composite with
unique and significantly improved mechanical properties
having a high ability to transfer the stresses from the
polymer matrix to the embedded particles. This allows
the silica/polymer nanocomposites to be used in a variety
of applications and industrial products successfully
replacing the classical materials. Although much work
has already been done on the silica/polymer nanocom-
posites, more research is needed to further understand
the complex filler-matrix relationship that would allow a
step forward and even enable a synthesis of the nano-
composites with controllable properties through tailoring
the interfacial interaction between silica and a polymer
matrix.
Acknowledgement
The author gratefully acknowledges the Slovenian
programme P2-0132, "Fizika in kemija povr{in kovin-
skih materialov" of the Slovenian Research Agency for
its support.
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Materiali in tehnologije / Materials and technology 47 (2013) 3, 285–293 293

M. NOSKO et al.: THE FATIGUE BEHAVIOUR OF ALUMINIUM FOAM
THE FATIGUE BEHAVIOUR OF ALUMINIUM FOAM
VEDENJE ALUMINIJEVIH PEN PRI PREIZKUSU UTRUJENOSTI
Martin Nosko, Franti{ek Siman~ík, Roman Florek
Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Ra~ianska 75, 831 02 Bratislava, Slovakia
ummsnoso@savba.sk
Prejem rokopisa – received: 2012-08-10; sprejem za objavo – accepted for publication: 2012-11-22
The aim of this work was to study the fatigue behaviour of aluminium foam during compression-compression cyclic loading
under an applied load. The reason was to estimate the amount of load for a fatigue life of 105 cycles with or without a small
amount of permanent plastic deformation determined to be less than 1 mm. The samples were subjected to cyclic loading under
various forces (proportion of level force estimated from a uni-axial compression test) in one given direction due to the
elimination of the foam anisotropy. Moreover, the fatigue behaviour of the aluminium foam under compression-compression
cyclic loading is described macroscopically. It was revealed that 50 % of the applied load estimated from the uni-axial
compression test is sufficiently low for a fatigue life 105 cycles in the case of a foam density of 0.211 ± 0.007 g cm–3.
Keywords: aluminium foam, fatigue life, level force, endurance limit
Cilj tega dela je bilo preu~evanje vedenja aluminijevih pen pri preizkusu utrujenosti s tla~no-cikli~no tla~nim obremenjevanjem
v odvisnosti od uporabljene obremenitve. Razlog je bil ugotoviti obremenitev za zdr`ljivost za utrujenost do 105 ciklov brez
majhne plasti~ne deformacije, manj kot 1 mm ter z njo. Vzorci so bili cikli~no obremenjeni z razli~nimi silami (proporcionalno
ravnote`ni sili, ugotovljeni pri enoosnem tla~nem preizkusu) v eni smeri, da bi izlo~ili vpliv anizotropije pene. Poleg tega je
makroskopsko predstavljeno vedenje aluminijeve pene pri preizkusu utrujenosti s tla~no-cikli~no tla~nim obremenjevanjem.
Ugotovljeno je, da je 50 % obremenitve, dobljene pri enoosnem tla~nem preizkusu, dovolj, da pena z gostoto 0,211 ± 0,007 g cm–3
pri utrujenostnem preizkusu zdr`i 105 ciklov.
Klju~ne besede: pena iz aluminija, zdr`ljivost za utrujenost, stopnja sile, maksimalna obremenitev
1 INTRODUCTION
Aluminium foam is known in industry as a good
candidate material for the cores of sandwich panels due
to its good stiffness-to-weight ratio.1–4 The effect of com-
position, manufacturing parameters, sound-absorption
properties and electrical behaviour on the foams’
properties was already studied for example in3,5–8. The
deformation mechanism of aluminium foam during
uni-axial compression and during uni-axial com-
pression-compression cyclic loading was studied in5,9–15
macroscopically and by using X-ray computed
tomography and surface strain mapping. These studies
were focused on determining the deformation modes and
revealing the structural character responsible for ore-
mature yielding. Our study is focused on contributing to
an estimate of the force at which the aluminium foam
can be used for 105 compression-compression cyclic
loadings.
2 EXPERIMENTAL
2.1 Material
Alporas® aluminium foamed block of composition Al
+ 1.5 % Ca + 1.6–3 % TiH2 2,8 was used for the study.
Since the aluminium foam block is anisotropic,8,16,17 the
middle area of the block, which is characterized by the
presence of the largest amount of non-uniformities
within the structure (the presence of the elongated pores,
pore agglomerates, fractured pore faces, microstructural
in-homogeneities, etc.) was chosen for an estimation of
the level force. The density of the represented area is
0.211 ± 0.007 g cm–3. Uni-axial compression-compres-
sion cyclic loadings were performed on a material testing
system device (MTS 810) using cubic samples of
dimension a = 45 mm. The deformation mechanism on
the macro-scale was revealed using a digital camera set
on an MTS device
2.2 Estimation of the level force
The uni-axial load-stress behaviour of the aluminium
foam has been studied elsewhere5–7 and is shown in
Figure 1a. It covers three regions: the elastic region at
low loads (includes the hardening region) ended by peak
stress, the long load-strain plateau wherein localized
plastic collapse propagates from one cell band to
another. After all bands collapse densification region of
rapid rise of load starts.
The level force (responsible for the plastic collapse of
the first cell band within the sample5–7) was estimated as
an average from 25 uni-axial compression tests accord-
ing to the DIN norm18 to 2200 N. It should be noted that
10 of the 25 samples showed a load peak well below
(1000 N and 1800 N, respectively) due to the occurrence
of large, elongated pores within the structure of the
represented area.
2.3 Compression-compression fatigue behaviour
To find out the relation between the ultimate pro-
portions of the level force (LF) and the fatigue behaviour
Materiali in tehnologije / Materials and technology 47 (2013) 3, 295–298 295
UDK 669.71:620.178.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)295(2013)
the samples were loaded by forces increases from 50 %
to 90 % LF with a step of approximately 5 % and 10 %.
The samples were subjected to compression-compres-
sion cyclic loading according to Figure 1b during the
period over 105 cycles. The fatigue life is defined as the
number of cycles corresponding to the onset of an abrupt
strain jump.
3 RESULTS AND DISCUSSION
3.1 Deformation mechanism
As presented in Figure 2a, the compression-com-
pression S-N curves during the cyclic loading are
comparable for applied fraction of LF, but the onset of
the yielding starts after different number of cycles. While
in the case of 90 % LF, it is achieved immediately after
the cyclic loading, in case of 69 % LF it starts after
30000 cycles and is characterized by changes in the
strain-rates. In the case of 63 % LF, it is achieved after
approximately 97000 cycles, which is almost the defined
endurance limit (105 cycles) – see detail in Figure 2b. It
suggests that the fatigue life (in case of a rigid material –
rapid coalescence and growth of cracks, in the case of
foamed material – evolution and subsequent collapse of
weak pore bands) is mainly affected by the proportion of
the loading force and the mechanism of fatigue damage
is delayed with a decrease of the fraction (%) LF.
The deformation mechanism during compression-
compression cyclic loading can be described through
S-N curves and is revealed macroscopically in Figure 3.
The principle of evolution of the plastic deformation and
its spreading within the sample during uni-axial com-
pression and cyclic loading is similar.9–12 Immediately
after loading in the elastic region, the uniform evolution
of the plastic hinges within the whole samples was
observed. The buckling and bending of the pore walls
(Figure 3a, b) together with the continuous evolution of
the plastic hinges follows as the deformation continues
into the "hardening region" and through the hardening
region.12 These features are responsible for the loss of
stiffness associated with a continuous shortening before
the abrupt strain jump presented in Figure 2b. The
abrupt strain jump is related to a macroscopically viewed
plastic collapse of a group of pores within the defor-
mation bands (Figure 3b, c) and causes changes in the
strain rate due to the softening effect after the collapse.12
Subsequently, in the vicinity of the already plastically
deformed group of pores the whole deformation band
M. NOSKO et al.: THE FATIGUE BEHAVIOUR OF ALUMINIUM FOAM
296 Materiali in tehnologije / Materials and technology 47 (2013) 3, 295–298
Figure 2: Compression-compression S-N curves; a) representation of
the different fatigue life with respect to the proportion of LF, b) detail
of the S-N curve
Slika 2: Tla~no-tla~ne S-N-krivulje; a) predstavitev razli~ne zdr`lji-
vosti pri utrujenostnem preizkusu v odvisnosti od LF, b) detajl
S-N-krivulje
Figure 1: a) Typical uni-axial compression load-strain curve with
definition of the level force (LF), b) description of the applied load
during the compression-compression cyclic loading
Slika 1: a) Zna~ilna krivulja sila – raztezek z opredeljeno stopnjo sile
(LF), b) videz uporabljene obremenitve med tla~no-cikli~no tla~nim
obremenjevanjem
closes up and the process is repeated within the other
region of the foamed sample (Figure 3c, d), also
accompanied by changes in the strain rate (Figure 2a,
No.11). However, the changes in the strain rates are
omitted if the loading is performed with a high
proportion of LF, as presented in Figure 2a No.14.
Vice-versa, no permanent deformation occurs if the
loading is performed by a small proportion of LF to
overcome the load of the plastic collapse of the pore
band, as seen in Figure 2a No.10.
3.2 Estimation of LF
As seen in Figure 4, the fatigue behaviour of the
foam is very sensitive to the proportion of LF = 2200 N.
While during loading at 50 % LF (Figure 4a), all the
samples achieved the endurance limit, the increment of
LF causes a decreased number of samples and is
responsible for the premature shortening of the samples
accompanied by changes in the strain rate, as presented
in Figures 4b to d. In case of 63 % LF, two of the six
samples achieved the endurance limit without premature
shortening; if 70 % of LF is applied, only one from three
samples reaches the limit, as presented in Figure 4c.
During cycling at 90 % LF (Figure 4d), all the samples
collapsed immediately after loading without changes in
the strain rate.
4 CONCLUSION
It was revealed that a 50 % of level force is sufficient
for uni-axial compression-compression cyclic loading to
achieve an endurance limit of 105 cycles with no
significant deformation, estimated to be less than 1mm.
In the elastic region, only the creation of plastic hinges
during the loading occurs, but in the region of hardening
the buckling and bending of the cell walls start, which
overcomes the reversible elastic deformation. The tests
were performed on samples chosen from an area with the
presence of the largest amount of non-uniformities
within the structure at 50 % LF. This suggests that 50 %
LF is the force that prevents the buckling and bending of
cell faces responsible for the plastic collapse of pore
bands. A force above 50 % LF causes continuous
damage to the foam before the endurance limit is
achieved.
Acknowledgements
The authors thank APVV-0647-10 Ultralight and the
Slovak Grant Agency - VEGA grant No. 2/0191/10 for
funding this work. Material support from Gleich®
GmbH is gratefully acknowledged. Infrastructure was
supported by CEKOMAT 2 with ID 26240120020. The
project is signed as OPVaV-2008/4.1/01-SORO.
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M. NOSKO et al.: THE FATIGUE BEHAVIOUR OF ALUMINIUM FOAM
Materiali in tehnologije / Materials and technology 47 (2013) 3, 295–298 297
Figure 4: S-N curves for compression-compression cyclic loading for
various LF proportion; a) 50 %, b) 63 %, c) 70 % and d) 90 %
Slika 4: S-N-krivulje za tla~no-cikli~no tla~no obremenjevanje pri
razli~nih LF-razmerjih; a) 50 %, b) 63 %, c) 70 % in d) 90 %
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M. NOSKO et al.: THE FATIGUE BEHAVIOUR OF ALUMINIUM FOAM
298 Materiali in tehnologije / Materials and technology 47 (2013) 3, 295–298
M. ÖBEKCAN et al.: ADDITION OF STRONTIUM TO AN Mg-3Sn ALLOY AND AN INVESTIGATION ...
ADDITION OF STRONTIUM TO AN Mg-3Sn ALLOY
AND AN INVESTIGATION OF ITS PROPERTIES
DODATEK STRONCIJA ZLITINI Mg-3Sn IN PREISKAVA NJENIH
LASTNOSTI
Mediha Öbekcan1, Aysun Ayday1, Hüseyin ªevýk2, Süleyman Can Kurnaz1
1Sakarya University, Engineering Faculty, Department of Metallurgical and Materials Engineering, Sakarya, Turkey
2Mersin University, Engineering Faculty, Department of Metallurgical and Materials Engineering, Mersin, Turkey
ckurnaz@sakarya.edu.tr
Prejem rokopisa – received: 2012-08-13; sprejem za objavo – accepted for publication: 2012-11-27
The effect of strontium additions in mass fractions of (0.05, 0.1, 0.2, 0.5 and 1) % to a magnesium alloy (Mg-3Sn) was
investigated in this work. The alloys were gravity cast under a controlled atmosphere. The mechanical properties and the
microstructures of the above alloys were examined and recorded. The results revealed that an addition of strontium to the above
alloy had significantly affected its microstructure. The X-ray diffraction results showed that in all of the obtained alloys the
main phases were 
-Mg and Mg2Sn, and that the strontium-based intermetallics were not detected. The hardness values
increased with the increasing strontium content. The highest yield strength, tensile strength and elongation were exhibited by
the Mg-%3Sn-%0.1Sr alloy.
Keywords: Mg-Sn alloy, Sr addition, mechanical properties
V tem delu je bil preiskovan u~inek dodatka masnih dele`ev stroncija (0,05; 0,1; 0,2; 0,5 in 1) % magnezijevi zlitini (Mg-3Sn).
Zlitine so bile gravitacijsko lite v kontrolirani atmosferi. Preiskane in ugotovljene so bile mehanske lastnosti in mikrostruture
navedenih zlitin. Rezultati so pokazali, da dodatek stroncija pri tej zlitini mo~no vpliva na mikrostrukturo. Rezultati rentgenske
difrakcije so pokazali, da so glavne faze 
-Mg, Mg2Sn v preiskovanih zlitinah, niso pa bile odkrite intermetalne faze na osnovi
stroncija. Vrednosti trdote so nara{~ale z nara{~ajo~o vsebnostjo stroncija. Najvi{ja meja plasti~nosti, najvi{ja trdnost in najve~ji
raztezek se je pokazal pri zlitini Mg-%3Sn-%0,1Sr.
Klju~ne besede: zlitina Mg-Sn, dodatek Sr, mehanske lastnosti
1 INTRODUCTION
There have been various investigations to identify the
properties of magnesium and its alloys with regard to
their use in industrial applications due to their low
density in comparison with the other commercial,
low-density alloys such as aluminium.1–5 Generally,
magnesium alloys are based on the Mg–Al system. For
example, aluminium-containing magnesium alloys will
have an Mg17Al12 compound, which adversely influences
the mechanical properties at high temperatures.6,7 It is
therefore important to add another alloying element to
reverse such an effect. Such elements are strontium (Sr),
calcium (Ca) and tin (Sn).8,9 In this work aluminium-free
magnesium alloys were used (Mg-3Sn).
2 EXPERIMENTAL WORK
The alloys were melted in a stainless-steel crucible
using an electric-resistance furnace facilitated with a
CO2-0.2SF6 atmosphere. Commercially pure magnesium,
tin and Mg-%20Sr were used in this case as shown in
Table 1. The melt was held at 760 °C for 10 min, then
stirred to ensure a homogeneous distribution of all the
alloying elements. The melts was then poured into a
preheated steel die at 270 °C. Cross-sections were taken
from similar areas of all the castings and ground down to
1200 grit using silicon carbide papers. The cross-sections
were then polished down to 1 μm samples using abrasive
diamond wheels. The specimens were chemically etched
using an acetic picric acid compound (5 ml of acetic
acid, 6 g of picric acid, 10 ml of distilled water, 100 ml
of ethanol), then examined using a JOEL scanning
electron microscope (SEM). The second lot of cross-
sections were taken from similar areas of all the castings.
The sections were ground using silicon carbide papers
before examining them with an X-ray diffraction [(XRD)
Rigaku D-Max 1000 X-ray diffractometer with Cu K

radiation] machine to identify their compounds and
phases.
Table 1: Chemical compositions of the investigated alloys (mass
fraction, w/%)
Tabela 1: Kemijska sestava preiskovanih zlitin (masni dele`i, w/%)
Alloys Composition Mg Sn Sr
1 Mg-3Sn 96.8 2.97 –
2 Mg-3Sn-0.05Sr 96.5 2.92 0.046
3 Mg-3Sn-0.1Sr 96.1 2.94 0.09
4 Mg-3Sn-0.2Sr 95.9 2.97 0.183
5 Mg-3Sn-0.5Sr 95.9 2.94 0.45
6 Mg-3Sn-1Sr 95.7 2.89 0.92
Brinell hardness tests were performed on all the
cross-sections using a diameter ball 2.5 mm with an
Materiali in tehnologije / Materials and technology 47 (2013) 3, 299–301 299
UDK 669.721.5:669.892:620.17 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)299(2013)
applied load of 31.25 kg. Tensile tests were carried out
using an Instron 3367 universal testing machine with a
fixed ram speed of 0.2 mm/s at ambient temperature.
3 RESULTS AND DISCUSSION
The XRD results shown in Figure 1 revealed that the
main phases were 
-Mg and Mg2Sn. The XRD results,
on the other hand, failed to show any Mg-Sr- and Sn-Sr-
based intermetallics in any of the tested castings. Here
the results are in agreement with the study made by
Hongmei Liu.9
The SEM images shown in Figure 2 revealed that the
microstructures of these alloys mainly consisted of the
primary 
-Mg surrounded by the boundary of the Mg2Sn
phase. It is very clear that the primary 
-Mg in the
strontium-free alloy 1 is much bigger than those seen in
alloys 2, 3, 4, 5 and 6. The intermetallics (Mg2Sn), on the
other hand, appear to be smaller and more nodular than
those of 2, 3, 4, 5 and 6. The increase in the strontium
content influenced the formation of a large Mg2Sn phase,
which also influenced the mechanical properties. Figure
3 shows the results of the energy-dispersive spectrometer
(EDS) obtained for alloy 6. The figure shows that stron-
tium was completely dissolved in the Mg2Sn inter-
metallic, which influenced its final shape.
The hardness increased and was directly proportional
to the increase in the strontium content as shown in
Figure 4. The tensile and the yield strengths increased
only up to alloy 3, then they decreased for alloys 4, 5 and
6. The elongation results showed a similar manner as
seen in the cases of the tensile and yield strengths. It is
clear from Figure 4 that the yield and tensile strengths
reached their maximum of about 78 MPa and 157 MPa
when 0.1 % Sr was used. The hardness, on the other
hand, reached a maximum of 40 BH when 1 % Sr was
used.
The reason for the increase in the tensile and yield
properties in the case of alloy 3 can be clearly attributed
to its microstructure. It is very clear from the microstruc-
ture that the longitudinal shape of the Mg2Sn inter-
metallic has influenced the increase in its tensile and
yield strengths.
M. ÖBEKCAN et al.: ADDITION OF STRONTIUM TO AN Mg-3Sn ALLOY AND AN INVESTIGATION ...
300 Materiali in tehnologije / Materials and technology 47 (2013) 3, 299–301
Mg Sn Sr O
1 79.266 9.935 0.961 9.837
2 80.628 8.260 0.699 10,412
3 84.764 4.952 2.634 7.649
4 75.648 11.145 0.438 12.770
5 71.894 11.235 2.091 14.780
6 98.733 1.267 – –
Figure 3: EDS analysis of alloy 6
Slika 3: EDS-analiza zlitine 6
Figure 1: XRD spectrums of selected alloys
Slika 1: XRD-spektri izbranih zlitin
Figure 2: SEM micrographs showing microstructures of: a) alloy 1, b) alloy 3 and c) alloy 6
Slika 2: SEM-posnetki mikrostrukture: a) zlitina 1, b) zlitina 3 in c) zlitina 6
4 CONCLUSIONS
1. The microstructures of all the above alloys consisted
of the 
-Mg and Mg2Sn phases.
2. The strontium addition modified the microstructural
shapes of the intermetallics.
3. The addition of Sr affected the UTS, yield and elon-
gation only to a certain extent.
4. The hardness increase was directly proportional to
the increase in strontium.
5 REFERENCES
1 A. Yu, S. Wang, N. Li, H. Hu, Journal of Materials Processing Tech-
nology, 191 (2007), 247–250
2 H. Hu, Journal of Materials Science, 33 (1998), 1579–1589
3 H. Q. Sun, Y. N. Shi, M. X. Zhang, Surface and Coating Technology,
202 (2008), 2859–2864
4 W. Huang, B. Hou, Y. Pang, Z. Zhou, Wear, 260 (2006), 1173–1178
5 T. A. Leil, N. Hort, W. Dietzel, C. Blawert, Y. Huang, K. U. Kainer,
K. P. Rao, Transactions of Nonferrous Metals Society of China, 19
(2009), 40–44
6 S. Kleiner, O. Beffort, A. Wahlen, P. J. Uggowitzer, Journal of Light
Metals, 2 (2002), 277–280
7 H. Liu, Y. Chen, H. Zhao, S. Wei, W. Gao, Journal of Alloys and
Compounds, 504 (2010), 345–350
8 S. Li, B. Tang, D. Zeng, Journal of Alloys and Compounds, 437
(2007), 317–332
9 H. Liu, Y. Chena, Y. Tang, S. Wei, G. Niu, Journal of Alloys and
Compounds, 440 (2007), 122–126
M. ÖBEKCAN et al.: ADDITION OF STRONTIUM TO AN Mg-3Sn ALLOY AND AN INVESTIGATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 299–301 301
Figure 4: Mechanical test results for the alloys: a) hardness and b)
strength and strain
Slika 4: Rezultati mehanskih preizkusov zlitin: a) trdota in b) trdnost
in deformacija

F. E. BAÞTAN et al.: THE EFFECT OF BINDER ON CHEMICALLY PRECIPITATED HYDROXYAPATITE ...
THE EFFECT OF BINDER ON CHEMICALLY
PRECIPITATED HYDROXYAPATITE DURING SPRAY
DRYING
VPLIV VEZIVA NA KEMIJSKO IZLO^ENI HIDROKSIAPATIT MED
ATOMIZACIJO
Fatih Erdem Baþtan, Ezgi Demiralp, Yýldýz Yaralý Özbek, Fatih Üstel
Sakarya University, Engineering Faculty, Department of Metallurgical and Metarials Engineering, Esentepe Campus, 54187 Sakarya, Turkey
yyarali@sakarya.edu.tr
Prejem rokopisa – received: 2012-08-17; sprejem za objavo – accepted for publication: 2012-11-22
The synthesis of appropriate calcium phosphate powders for thermal-spraying applications is a fundamental, crucial stage in the
production of bioceramical coatings coupled with the desired characteristics. The performance, lifespan and quality of the
resulting biological coating in-vivo is largely dependent on the coating morphology, phase composition, particle size and the
crystallites of the spray powders. In order to achieve very reliable coatings from thermal-spray processes, spherical powders of a
specified size distribution are recommended. The aim of this work was to produce hydroxyapatite powder with a chemical
precipitation method and to reshape it in a spray dryer and investigate the effect of binder on the powder structure to provide an
insight into the preparation and characterization aspect of HA powders using the spray-drying process. Ethanol, pure water and
polyvinilalcohol (PVA) + ethanol were used as the binder. Different temperatures were applied in the spray dryer. Then, the
precipitated, spray-dried powders were examined for morphology. Scanning electron microscopy (SEM), X-ray diffraction
(XRD), (EDX) and ICP were used to characterize the specimen powders.
Keywords: hydroxyapatite, chemical precipitation, spray dryer, ICP (Inductively Coupled Plasma)
Sinteza primernega prahu kalcijevega sulfata za termi~no napr{evanje, povezana z `elenimi lastnostmi, je osnovna in klju~na
faza pri izdelavi biokerami~nih prevlek. Uspe{nost, zdr`ljivost in kvaliteta biolo{ke prevleke v `ivo je mo~no odvisna od
morfologije prevleke, fazne sestave, velikosti delcev in kristalnih zrn napr{enega prahu. Za zagotovitev zelo zanesljivega
premaza se priporo~a uporaba prahu z okroglimi delci dolo~ene porazdelitve velikosti zrn. Cilj tega dela je izdelati prah
hidroksiapatita (HA) z metodo kemijskega izlo~anja, s preoblikovanjem z atomizacijo in preiskati u~inek veziva na strukturo
prahu, da bi dobili vpogled v na~ine priprave in karakterizacijo HA-prahov z atomizacijo. Kot veziva so bili uporabljeni etanol,
~ista voda in polivinil alkohol, (PVA) + etanol. Pri atomizaciji so bile uporabljene razli~ne temperature. Nato je bila pregledana
morfologija atomiziranega prahu. Karakterizacija vzorcev prahov je bila izvr{ena z vrsti~no elektronsko mikroskopijo (SEM),
rentgensko difrakcijo (XRD), energijsko disperzijsko rentgensko spektroskopijo (EDX) in induktivno sklopljeno plazmo (ICP).
Klju~ne besede: hidroksiapatit, kemijsko izlo~anje, atomizacija, induktivno sklopljena plazma ICP
1 INTRODUCTION
Bone is formed by collagen fibres and hydroxyapatite
natural bone tissue can be considered as a composite
consisting of a mineralized collagen matrix.1 Hydro-
xyapatite (HAp) Ca10(PO4)6(OH)2 and other related
calcium phosphate minerals have been evaluated as
implant materials for many years due to their good
biocompatibility and bioactivity as well as their simi-
larity with the inorganic components of the hard tissues
in natural bones. Their Ca/P ratio of 1.5–2.0 makes them
an excellent choice for most dental and orthopedic appli-
cations in the form of bioceramic coatings. Moreover,
HA has been used as a biological chromatography
support in protein purification and DNA isolation. Also,
HA is currently used for the fraction and purification of a
wide variety of biological molecules, such as subclasses
of enzymes, antibodies fragments and nucleic acids.2–6
Several methods, such as precipitation, solid-state
synthesis, hydrolysis, wet chemical, hydrothermal and
sol-gel methods have been used to prepare synthetic
HAp. The synthetic HAp is used for coating in medical
applications.
The HAp coating produced by plasma-spraying tech-
nology combines the mechanical advantages of a metal
substrate with the excellent biological properties of HAp.
Some important factors are the particle size, particle size
distribution and particle morphology, which affect the
lifetime and quality of the resulting biological coating.7,8
These important factors determine the flow charac-
teristics in the powder-feeding systems and the melting
behavior in the plasma jet.2,3
The hydroxyapatite powder size is very important for
a thermal-spray coating system. Therefore, we have to
increase the size of the powder for a good flow rate. The
spray-dryer system was used to adjust the particle size.
The spray-drying method is a kind of granule production
technique. The advantages of this method are very
simple and the particle size can be controlled quite
easily. The process parameters are the slurry concentra-
tion, the compressed-air flow rate and the liquid flow
rate, which are affected by the specific surface area and
size distribution of the final products.4 The morphology
of the powders is generally spherical, the other morpho-
logies are, for example, mushroom-like,9 doughnut-like,9
hollow structures10 etc.
Materiali in tehnologije / Materials and technology 47 (2013) 3, 303–306 303
UDK 66.09 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)303(2013)
In this study we fabricated HAp powders with the
chemical precipitation method and the produced powders
were granulated as spherical powders using the spray-
drying method by controlling the process parameters.
Different binders (pure water, PVA (polyvinyl alcohol)
and ethanol) were used for the slurry. The resultant
spherical powders were investigated to see the effect of
binders on the powder properties (morphology and
particle size distribution). The final spherical powders
were prepared as bulk materials and sintered. Scanning
electron microscopy (SEM), X-ray diffraction (XRD)
and inductively coupled plasma (ICP) were used to
characterize the powders and the bulk materials.
2 MATERIALS AND METHOD
HAp particles were synthesized by a chemical pre-
cipitation method, with calcium nitrate tetrahydrate
(Ca(NO3)2·4H2O) as the calcium source, phosphoric acid
(H3PO4) as the phosphorous source, and ammonium
hydroxide(NH4OH) as the pH regulator. (Ca(NO3)2 ·
4H2O) and (H3PO4) were separately dissolved in distilled
water continuously for 30 min. The dissolved solutions
were mixed together and added (NH4OH) to obtain the
initial pH values of the reaction solutions as 11.00. The
mixture was stirred at a speed of 250 r/min. The resulting
suspension was aged for 24 h at room temperature and
then filtered. The product was washed with water to
remove the residual impurities. The precipitated powders
were dried at 105 °C to remove the undesired impurities.
The dried powders were mixed with pure water, and PVA
and ethanol were used as a binder to obtain a slurry for
the spray dryer. Inlet temperatures of 175 °C, 190 °C and
200 °C and a 1.5 bar pressure were chosen for the spray
drying. The final product was shaped as a bulk material
and sintered at 1050 °C for 1 h.
3 RESULTS AND DISCUSSION
The XRD result of the hydroxyapatite after the sin-
tering process is shown in Figure 1. The peaks are sharp
and match with the reference hydroxyapatite peaks. The
powder has a crystalline structure. These results revealed
that hydroxyapatite with a chemical precipitation method
could be produced. Also, the composition of the powder
was given in Table 1 (i.e., the results of the ICP). It was
shown that the powder not only has Ca and P elements,
but also has Fe, Mg and Zn.
Figue 2a shows that the powders had an irregular,
and an angular shape distribution, and also range widely.
The SEM micrographs of the powder in Figures 2b
and 2c revealed that the powders produced by spray
drying had a less spherical and porous microstructure.
The high degree of porosity could be due to the
elimination of the binder that was used in the binding
and agglomeration of the spray-dried powder. Moisture
and gases were also released and eliminated as a result of
spray drying at an elevated temperature of 200 °C.11
Also, these porosities appear as a small surface
depression (Figures 2c and 2d), which may be mini-
mized and the material restored to a denser structure
through calcination or sintering. Visible colour changes
were seen in powder. The colour changes were due to the
presence of manganese ions or other transition-metal
elements located in the crystal lattice structure. Although
they may not have any significant effect on the biocom-
F. E. BAÞTAN et al.: THE EFFECT OF BINDER ON CHEMICALLY PRECIPITATED HYDROXYAPATITE ...
304 Materiali in tehnologije / Materials and technology 47 (2013) 3, 303–306
Figure 2: a) Hydroxyapatite microstructure before spray drying, b)
after spray drying (175 °C, 1.5 bar, ethanol), c) (200 °C, 1.5 bar, etha-
nol), d) (175 °C, 1.5 bar, ethanol + PVA)
Slika 2: a) Mikrostruktura hidroksiapatita pred atomizacijo, b) po ato-
mizaciji (175 °C, 1,5 bar, etanol), c) (200 °C, 1,5 bar, etanol), d) 175
°C, 1,5 bar, etanol + PVA)
Figure 1: XRD peaks after sintering
Slika 1: XRD-spekter po sintranju
Table 1: ICP analysis result
Tabela 1: Rezultati ICP-analize
Composition Amount in mass fractions, w/%
PO4 61.05
Ca 38.50
Fe 0.0030
Mg 0.12
Zn 0.0068
patibility of HA, the consumer acceptance should be
duly considered.12
When the binder is ethanol for drying, the optimal
parameters are 175 °C inlet temperature and 1.5 bar pres-
sure, for pure water the optimal parameters are 175 °C
inlet temperature and 1.5 bar pressure.13 Spherical
particles are seen in Figure 2b. The powders have a
porous structure because of the early evaporation of the
ethanol. A porous and hollow structure with the slow
diffusion of solute and a quick solvent evaporation were
obtained.14 Irregular particles are seen in Figure 2c.
Increasing the inlet temperature results in quick evapo-
ration of the moisture, but a high temperature may cause
chemical/physical distortion.15 The spherical and porous
particles are seen in Figure 3. Ethanol + PVA were
chosen as binders. The results are same as those obtained
from the ethanol-added samples. The PVA affected only
the particle size. After the particle size analysis, it was
seen that the particle size increases with increasing PVA
addition for a good binding. The average particle sizes
were 27 μm and 41 μm for the ethanol and ethanol +
PVA, respectively.
It was shown in Figures 3a and 3b that spherical
particles were obtained at both 175 °C and 190 °C, but it
seems that particles have moisture because of the
insufficient inlet temperature for drying in Figure 3a.
Increasing the temperature made the particles dry.
Particles have less porosity when using pure water for
the binding. Because the pure water’s evaporation tempe-
rature is higher than the ethanol’s, the binder holds
together all particles during the process.15
4 CONCLUSIONS
Hydroxyapatite powders could be produced by a
chemical precipitation method and reshaped with a spray
dryer. The change in the binder impacted on the
spray-drying parameters. The optimal inlet temperatures
are 175 °C for ethanol and 190 °C for pure water. The
type of binder was affected by the particle structure. A
volatile binder resulted in a lower particle density and
the particles had more porosity. The binder holds
together all the particles, and increasing the amount of
the binder (like PVA) increases the particle size.
Spray-dried powder with the correct particle size is
converted to flame spheroidized powder so as to improve
the microstructural characteristics and the stability of the
powder. A spherical geometry is very desirable for
enhanced flowability and deposition consistency, which
would eventually give rise to high-quality bioceramic
coatings.
5 REFERENCES
1 F. Chen, Z. C. Wang, C. J. Lin, Materials Letters, 57 (2002),
858–861
2 O. Kikuo, A. I. Mikrajuddin, L. Wuled, I. Ferry, Preparation of func-
tional nanostructured particles by spray drying, Advanced Powder
Technol., 17 (2006) 6, 587–611
3 S. W. K. Kweh, K. A. Khor, P. Cheang, An in vitro investigation of
plasma sprayed hydroxyapatite (HA) coatings produced with flame-
spheroidized feedstock, Biomaterials, 23 (2002) 3, 775–785
4 E. S. Thian, K. A. Khor, N. H Loh, S. B. To, Processing of
HA-coated Ti–6Al–4V by a ceramic slurry approach: an in vitro
study, Biomaterials, 22 (2001) 11, 1225–1232
5 A. Afshar, M. Ghorbani, N. Ehsani, M. R. Saeri, C. C. Sorrell, Some
important factors in the wet precipitation process of hydroxyapatite,
Materials & Design, 24 (2003) 3, 197–202
6 E. Lugscheider, M. Knepper, K. A. Gross, J. Thermal Spray Tech., 1
(1992) 3, 215–222
7 S. W. K. Kweh, K. A. Khor, P. Cheang, Journal of Materials Process-
ing Technology, 89–90 (1999), 373–377
8 A. J. Wang, Y. P. Lu, R. F. Zhu, S. T. Li, X. L. Ma, Powder Tech-
nology, 191 (2009), 1–6
9 H. Liang, K. Shinohara, H. Minoshima, K. Matsushima, Analysis of
constant rate period of spray drying of slurry, Chem. Eng. Sci., 56
(2001), 2205–2213
F. E. BAÞTAN et al.: THE EFFECT OF BINDER ON CHEMICALLY PRECIPITATED HYDROXYAPATITE ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 303–306 305
Figure 3: a) After spray drying (175 °C, 1.5 bar, pure water), b) (190
°C, 1.5 bar, pure water)
Slika 3: a) Po atomizaciji (175 °C, 1,5 bar, ~ista voda), b) (190 °C, 1,5
bar, ~ista voda)
10 R. Sun, Y. Lu, K. Chen, Materials Science and Engineering C, 29
(2009), 1088–1092
11 F. E. Bastan, Kimyasal Çöktürme Yöntemiyle Gümüþ Ýlaveli Hidro-
ksiapatit Üretimi ve Spray Dryer ile Þekillendirilmesi, Master thesis,
Sakarya university, 2012
12 S. W. K. Kweh, K. A. Khor, P. Cheang, The production and characte-
rization of hydroxyapatite (HA) powders, Journal of Materials Pro-
cessing Technology, 89–90 (1999), 373–377
13 T. Baþargan, Püskürtmeli Kurutucuda Hidroksiapatit-Polimer Malze-
melerin Hazýrlanmasý, 2010
14 K. Mehsana, Spray Drying Technology, An Overview, 2009
15 R. Chumnanklang, T. Panyathanmaporn, K. Sitthiseripratip, J. Su-
wanprateeb, Freeform Fabrication of Hydroxyapatite via Three
Dimensional Printing, Materials Science and Engineering, 27 (2007)
4, 914–921
306 Materiali in tehnologije / Materials and technology 47 (2013) 3, 303–306
F. E. BAÞTAN et al.: THE EFFECT OF BINDER ON CHEMICALLY PRECIPITATED HYDROXYAPATITE ...
L. C. KUMRUOGLU, A. ÖZEL: PLASMA ELECTROLYTIC SATURATION OF 316 L STAINLESS STEEL ...
PLASMA ELECTROLYTIC SATURATION OF 316 L
STAINLESS STEEL IN AN AQUEOUS ELECTROLYTE
CONTAINING UREA AND AMMONIUM NITRATE
PLAZEMSKO ELEKTROLITSKO NASI^ENJE
NERJAVNEGA JEKLA 316 L V VODNEM ELEKTROLITU
S SE^NINO IN AMONIJEVIM NITRATOM
Levent Cenk Kumruoglu, Ahmet Özel
University of Sakarya, Engineering Faculty, Metallurgy and Materials Department, Esentepe-Sakarya, Turkey
lkumruoglu@sakarya.edu.tr
Prejem rokopisa – received: 2012-08-30; sprejem za objavo – accepted for publication: 2012-11-13
Plasma electrolytic saturation (PES) is an environmentally friendly electrochemical process that allows altering the surface
chemistry and grain size of metallic substrates, negatively biased in an aqueous electrolyte containing ionic species such as N, O
and C. Wear- and corrosion-resistive nanocrystalline layers consisting of carbides, nitro carbides, borides and nitro-carbon
oxides could be fabricated with PES in short treatment durations. In this study, PES was performed on 316 L stainless steel in an
electrolyte containing urea and ammonium nitrate with several treatment durations from 5 s to 30 min. The surface morphology,
topography and microstructure were investigated with X-Ray diffraction, optical microscopy, scanning electron microscopy
with energy dispersive spectroscopy, a surface profilometer and microhardness testing. The wear and friction properties were
evaluated using a ball-on-plate, linear, reciprocating wear test at 1 N to 3 N applied loads with an alumina ball against both the
treated and untreated substrates. It was found that PES can increase the wear resistance and mechanical properties of 316 L
stainless steel.
Keywords: plasma electrolysis, saturation – diffusion, nitriding, wear, 316 L
Plazemsko elektrolitsko nasi~enje (PES) je okolju prijazen elektrokemijski postopek, ki omogo~a spreminjanje povr{inske
kemije in velikost zrn v kovinski podlagi, ki je neustrezna v vodnem elektrolitu, ki vsebuje ione N, O in C. V kratkem ~asu je
mogo~e s PES izdelati obrabno in korozijsko obstojne nanokristalne plasti iz karbidov, nitrokarbidov, botidov in
nitro-karbo-oksidov. V tej {tudiji je bil PES uporabljen pri nerjavnem jeklu 316 L v elektrolitu, ki je vseboval se~nino in
amonijev nitrat in so bili ~asi obdelave od 5 s do 30 min. Preiskovana je bila morfologija povr{ine, topografija in mikrostruktura
z rentgensko difrakcijo, svetlobno mikroskopijo, vrsti~no elektronsko mikroskopijo, energijsko disperzijsko spektroskopijo, s
povr{inskim profilometrom in meritvijo trdote. Obraba in torne lastnosti so bile ocenjene z metodo kroglica na plo{~i z
linearnim izmeni~nim obrabnim preizkusom pri obte`bi od 1 N do 3 N s kroglico iz Al2O3 na obdelani in neobdelani povr{ini.
Ugotovljeno je bilo, da PES lahko pove~a odpornost proti obrabi in izbolj{a mehanske lastnosti nerjavnega jekla 316 L.
Klju~ne besede: plazemska elektroliza, nasi~enje – difuzija, nitriranje, obraba, 316 L
1 INTRODUCTION
To improve and alter the surface properties of
metallic substrates with a plasma-electrolytic-saturation
process,1–7 interstitial atoms like C, N, O and species are
diffused into the surface of a biased substrate from an
aqueous medium containing interstitial atoms under the
appropriate plasma conditions like specific temperature,
potential, duration.1–5 Plasma electrolysis is a special
thermo-chemical-mechanical process employing electro-
lysis in an aqueous solution under particular conditions,
for instance specific potential, current, electrolyte and
durations.1–7 Plasma electrolysis is a complex process,
which couples physical metallurgy and electrochemical
events, such as heating a work piece in a cathodic regi-
me2,4 (gas liberation, spark ignition, continuous plasma
envelope, and arcing regime), where phase transforma-
tions and deformations occur simultaneously.
For the PES process, the selection of an electrolyte is
relatively simple; for example, for the nitrocarbusing the
electrolyte is composed of C/N-containing organic com-
pounds in a conductive solution, e.g., KCl or Na2CO3. In
contrast to conventional pack and liquid-based treat-
ments using cyanide, various ecologically friendly orga-
nic compounds can be used to provide, with plasma
thermal decomposition, desirable carbon/nitrogen ions
and/or atoms for the treatment. PES can have a consi-
derable strengthening effect on steel substrates. Even
though a wide range of materials has been treated using a
plasma-electrolytic-saturation technique,2 very few
studies exist in the literature that report on the wear
resistance of oxy-nitrided or oxy-carbon-nitrided 316 L
stainless steel. To evaluate the saturation phenomena and
wear behaviour of the oxy-nitrided layer, different elec-
trolytes, process durations and dry, reciprocating. sliding
tests have been conducted. The modified phases on the
surface were investigated with XRD (Rigaku) applying
the standard 2-scan from 20 to 90 degrees and cross-
sections of saturated layers were revealed with optical
microscopy (Nikon) and scanning electron microscopy
(JEOL 6600) with energy dispersive spectroscopy. The
Materiali in tehnologije / Materials and technology 47 (2013) 3, 307–310 307
UDK 621.357:669.058 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)307(2013)
surface roughness and microhardness of the samples
were measured using a surface profilometer (Tencor P6)
and a microhardness tester (Shimadzu HMV).
2 EXPERIMENTAL STUDIES
The testing coupons (10 mm × 20 mm × 1.5 mm)
made from 316 L stainless steel were used as the sub-
strate materials. The nominal composition of a substrate
in mass fractions was Cr 18 %, Mn 1.6 %, Ni 10 %, C
0.02 %, Mo 2.1 %, N < 0.02 %, Fe balance. Prior to
PES, the substrates were polished to Ra  0.05 μm from
240 to 1200 grit with the emery paper and then with the
alumina solution. The PES was carried out by using 24
kW DC power supply and in a specially designed
instrumented rig shown in Figure 1. Two types of elec-
trolyte at 25 °C were chosen as given in the following
formulas, m(E1) = 4 kg urea, 40 g KOH and 4 l di-water
(cond. 8 mS) and m(E2) = 4 kg ammonium nitrate, 4 l
di-water, 20 g KOH (cond. 180 mS). The substrates were
biased negatively and treated for 5 s and (1, 15 and 30)
min at the potentials varying from 100 V (for E2) to 300
V (for E1). The breaking potentials were 550 V for E1
and 350 V for E2. The electrolyte was pumped from the
container to the bottom surface of the samples and then
the potential was increased manually to fix the plasma
envelope. The treated surfaces were tested using the
following parameters: the normal loads of 1 N and 3 N,
the track length of 10 mm, the frequency of 5 Hz, the
sliding speed of 0.1 m s–1 and the sliding distances of
100 m and 1200 m. The average value of the friction
coefficient was measured in both reciprocating directions
of the bearing ball. The CSM tribometer was used for the
wear tests. Before and after each test, both the sample
and the alumina ball (10 mm) were cleaned ultrasoni-
cally in acetone and dried with hot air.
3 RESULTS AND DISCUSSION
Recent studies have revealed that plasma saturation
can produce compound nitrides and oxide phases, e.g.,
(Fe,Cr)2O4, Fe3O4+FeN, CrN, (N).3–6 The XRD patterns
of the samples treated with E1 and E2 are given in
Figures 2a and b. SEM and OM studies revealed that the
surfaces consist of Fe, Cr2O4, Fe3O4, FeN0.076 and a N
saturated layer that is 3 μm to 40 μm thick. In addition to
the mixture phases, Fe2O3 was detected in the sample
saturated in urea for 5 s. According to the Fe-O binary
system and the phase-stability diagram, Fe2O3 and Fe3O4
can be formed at 400 °C, depending on the oxygen con-
centration, and FeO is not stable below 570 °C. In the
Fe-O system, a larger amount of oxygen causes a forma-
tion of hematite at high temperatures. However, in the H2
atmosphere, Fe2O3 may be reduced to Fe3O4.8 In plasma
electrolysis hydrogen can be produced at the cathode.3
Hence, longer plasma durations may cause the formation
of larger and thicker amounts of Fe3O4 and a broadening
of the XRD lines. SEM-EDS analyses proved the nitro-
gen, oxygen and carbon diffusions for both plasma pro-
cesses (Figure 3). The microhardness testing was
performed on the saturated layers. The maximum value
of 535 HV was observed for the PES using the urea
electrolyte for 30 min. However, the hardness testing
L. C. KUMRUOGLU, A. ÖZEL: PLASMA ELECTROLYTIC SATURATION OF 316 L STAINLESS STEEL ...
308 Materiali in tehnologije / Materials and technology 47 (2013) 3, 307–310
Figure 2: XRD of substrates saturated with electrolytic plasma with:
a) urea and b) with ammonium nitrate
Slika 2: XRD podlage, nasi~ene z elektrolitsko plazmo: a) s se~nino
in b) z amonijevim nitratom
Figure 1: Schematic drawing and picture of the PES setup
Slika 1: Shematski prikaz PES-sestava
cannot be performed on the PES substrates using an
ammonium nitrate electrolyte because of the rough and
thin hardened layer. On the other hand, the surface hard-
ness was measured as 550 HV using light indentation
loads such as 10 g (Figure 4). Recently, the post-oxida-
tion treatment was studied on nitrided layers.9 The
studies suggest that a thin Fe3O4 layer improves the wear
resistance of a nitrided layer.9 PES may be described as a
single-step compound diffusion process that enables both
nitriding and post oxidation. Nitriding can be performed
in a relatively low temperature range of 400–600 °C.2,3 In
PES, the surface of the substrate may reach this tempe-
rature in a few seconds.3 It is possible to heat a substrate
above the nitriding temperatures; if the temperatures
increase the oxidation may occur on the nitrided layer
because of the oxygen atoms in the electrolyte. The wear
behavior and the friction coefficient of the saturated
layers sliding against the alumina ball 10 mm under a
load of 3 N is shown in Figures 5a and b, respectively. It
can be seen that the wear resistance of the substrate
saturated for (1, 15 and 30) min in the E2 electrolyte is
much greater than that of the untreated substrate. The
untreated sample was marked with the symbol "0". Also,
there are a small drop for the PES 1 min and sharp drops
for the PESs 15 min and 30 min in the friction coeffi-
cient over the sliding distance 1200 m. The Fe3O4 layer
on the saturated steel may cause this kind of friction
behavior.9 The wear resistance of the substrate treated
with the urea electrolyte was improved after the PES 15
min. On the other hand, a good result was observed for
the PES 3 min. This may be related to both an increase
in the surface hardness and the thickness of the oxy-
nitrided layer. The thin layers (treated for 5 s) were
probably delaminated at the onset of the wear test
because there was no significant change in the friction
coefficients of both electrolyte groups. Below the sliding
distance 100 m the wear resistance of the PES layers
treated for 5 s in the ammonium nitrate electrolyte was
2.5 times better than that of the untreated substrate,
while for the sample treated for 30 min the wear
resistance was 10 times higher than that of the untreated
one and the worn track was very thin and hardly
L. C. KUMRUOGLU, A. ÖZEL: PLASMA ELECTROLYTIC SATURATION OF 316 L STAINLESS STEEL ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 307–310 309
1 w/% 2 w/% 3 w/% 4 w/%
C 21.41 C 1.007 C 3.120 C 3.385
N 0.000 N 0.678 N 0.498 N 0.000
O 41.10 O 38.96 O 19.400 O 2.017
Fe 35.06 Cr 9.472 Cr 22.418 Cr 18.07
Ni 1.898 Fe 47.94 Fe 44.217 Fe 68.78
Mo 0.519 Ni 1.579 Ni 9.984 Ni 7.540
100.0 Mo 0.354 Mo 0.363 Mo 0.199
100.0 100.0 100.0
Figure 3: Cross-section SEM view and EDS analysis of a saturation
layer (1 min with E1)
Slika 3: Prerez (SEM) in EDS-analiza nasi~ene plasti (1 min z E1)
Figure 5: a) Wear properties of PES layer for 3 N and 1200 m sliding
distance and b) the friction coefficient
Slika 5: a) Obrabne lastnosti PES-plasti pri 3 N in 1200 m drsne poti
in b) torni koeficienti
Figure 4: Cross-sectional microhardness profile of saturated samples
Slika 4: Mikrotrdota na pre~nem prerezu nasi~enih vzorcev
observed (Figure 6a). On the other hand, the wear-test
results for the samples treated with the urea solution
were slightly better than the results for the ammonium
nitrate electrolyte (Figure 6b).
4 CONCLUSIONS
The electrolytes containing urea and ammonium
nitrate can be used for the surface saturation of 316 L
stainless steel. The PES layers consist predominantly of
the FeCr2O4, FeN and Fe3O4 phases. In short PES dura-
tions, Fe2O3 was observed with the other phases because
of the low substrate temperature. At the beginning of
PES, the expanded austenite can be formed with an
atomic diffusion; however, an increase in the temperature
of the substrate gives rise to a formation of metal (nitro)
oxides. The H2 gas at the cathode causes a reduction of
hematite; so a stable mix of the FeCr2O4, FeN0,076 and
Fe3O4 phases can be formed on the surface. These phases
increase the hardness and reduce the friction coefficient,
thus increasing the wear resistance. There was no scaling
or delaminating of the oxide layers under the all-mecha-
nical tests.
The value of the breaking potential depended on the
conductivity of the electrolyte. The breaking potentials
of the electrolyte were 550–600 V for the conductivity of
8 mS (E1) and 300–350 V for the conductivity of 180
mS (E2).
The layer thickness increases with the treatment time,
resulting in an associated increase in the surface hard-
ness.
5 REFERENCES
1 M. Aliofkhazrae, A. S. Rouhaghdam, T. Shahrabi, Journal of Alloys
and Compounds, 460 (2008), 614–618
2 A. L. Yerokhin et al., Surface and Coatings Technology, 122 (1999)
2–3, 73–93
3 C. Cionea, PhD thesis, The University of Texas At Arlington, 2010
4 T. Paulmier, J. M. Bell, P. M. Fredericks, Thin Solid Films, 515
(2007), 2926–2934
5 P. Taheri, C. Dehghanian, Transaction B: Mechanical Engineering,
16 (2009) 1, 87–91
6 L. C. Kumruoðlu, A. Özel, Materials and Manufacturing Processes,
25 (2010) 9, 923–931
7 L. C. Kumruoglu, D. A. Becerik, A. Ozel, A. Mimaroglu, Materials
and Manufacturing Processes, 24 (2009), 781
8 Lin et. al., Thermochimica Acta, 400 (2003), 61–67
9 M. L. Fares, K. Chaoui, J. Le Coze, Surf. Interface Analysis, 41
(2009), 4549–4559
L. C. KUMRUOGLU, A. ÖZEL: PLASMA ELECTROLYTIC SATURATION OF 316 L STAINLESS STEEL ...
310 Materiali in tehnologije / Materials and technology 47 (2013) 3, 307–310
Figure 6: Wear properties of PES layer for 3 N and 100 m sliding
distance for both electrolytes
Slika 6: Obrabne lastnosti PES-plasti pri 3 N in 100 m drsne poti za
oba elektrolita
J. STETINA et al.: MINIMIZATION OF SURFACE DEFECTS BY INCREASING THE SURFACE TEMPERATURE ...
MINIMIZATION OF SURFACE DEFECTS BY
INCREASING THE SURFACE TEMPERATURE DURING
THE STRAIGHTENING OF A CONTINUOUSLY CAST
SLAB
ZMANJ[EVANJE POVR[INSKIH NAPAK Z ZVI[ANJEM
TEMPERATURE POVR[INE KONTINUIRNO ULITEGA SLABA
MED RAVNANJEM
Josef Stetina, Tomá{ Mauder, Lubomir Klimes, Frantisek Kavicka
Brno University of Technology, Technicka 2, 616 69 Brno, Czech Republic
stetina@fme.vutbr.cz
Prejem rokopisa – received: 2012-08-31; sprejem za objavo – accepted for publication: 2012-10-23
Surface temperatures of cast slabs on small-radius segments as well as on the unbent areas belong to the parameters that affect
the surface quality of continuously cast slabs. Older machines for continuous casting were designed with regard to the quantity
(the amount of cast slabs) rather than the quality of the production. Therefore, an adaptation of the secondary cooling is required
in order to obtain the desired surface temperatures. The modification consists of a dynamic control of the secondary cooling,
surface-temperature monitoring by means of a numerical model of the temperature field as well as a prospective replacement of
the cooling nozzles. In order to optimize and control the secondary cooling, characteristics of the nozzles, especially the
influences of the water-flow rate, air pressure, casting speed, surface temperatures and heat-transfer coefficient under the
nozzles have to be known. Moreover, the heat-transfer coefficient can also be influenced by the age of the nozzles. The paper
deals with the relationships between these influences and their impacts on the temperature field of a cast slab. The results are
presented for the 1530 mm × 250 mm slabs that are cast in Evraz Vítkovice Steel where the main author’s dynamic, 3D
solidification model is used, in its off-line version, to control the production interface. The results can be used for the
preparation of a real casting process.
Keywords: optimization of the temperature field, surface temperature of a slab, characteristics of nozzles, continuous casting
Temperatura povr{ine ulitega slaba pri segmentih z majhnim radiusom, kot tudi na neukrivljeni povr{ini, spada k parametrom, ki
vplivajo na kvaliteto povr{ine kontinuirno ulitega slaba. Starej{e naprave za kontinuirno ulivanje slabov so bile pripravljene bolj
za ve~jo zmogljivost kot pa za kvaliteto. Zato je potrebna prilagoditev sekundarnega ohlajanja, da se zagotovi doseganje `elene
temperature povr{ine. Prilagoditev sestoji iz dinami~ne kontrole sekundarnega hlajenja, kontrole temperature povr{ine z
numeri~nim modelom temperaturnega polja, kot tudi morebitna zamenjava hladilnih {ob. Optimiranje in kontrola sekundarnega
hlajenja je mogo~a s poznanjem zna~ilnosti {ob in {e posebno vpliva hitrosti pretoka vode, tlaka zraka, hitrosti ulivanja,
temperature povr{ine in koeficienta prenosa toplote pod {obami. Poleg tega na koeficient prenosa toplote lahko vpliva tudi
starost {ob. Ta ~lanek obravnava odnos med na{tetimi vplivi in njihov u~inek na temperaturno polje ulitega slaba. Predstavljeni
so rezultati za ulit slab 1530 mm × 250 mm. Avtorjev dinami~ni 3D-model strjevanja se uporablja za kontrolo vmesnika pri
proizvodnji in te~e v off-line-verziji. Rezultati se lahko uporabijo kot pripravljalno orodje za realni postopek ulivanja.
Klju~ne besede: optimiranje temperaturnega polja, temperatura povr{ine slaba, zna~ilnosti {ob, kontinuirno ulivanje
1 INTRODUCTION
The presented in-house model of the transient-
temperature field of the blank from a slab caster (Figure
1) is unique as, in addition to being entirely 3D, it can
work in real time. The numerical model covers the
temperature field of the complete length of the blank
(i.e., from the meniscus inside the mould all the way
down to the cutting torch) with up to one million nodes.1
The concasting machine (caster) for the casting of
slabs (Figure 1) has the secondary-cooling zone sub-
divided into thirteen sections due to the convection of a
greater amount of heat from the voluminous slab casting.
The first section engages the water nozzles from all sides
of a slab. The remaining twelve sections engage air-mist
cooling nozzles, positioned only on the upper and lower
sides of the concasting. It is therefore very important to
determine the correct boundary conditions for a numeri-
Materiali in tehnologije / Materials and technology 47 (2013) 3, 311–316 311
UDK 621.74.047:519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)311(2013)
Figure 1: Radial caster and positions of the nozzles along the slab
caster in 13 individual zones
Slika 1: Livni stroj z radijem in pozicija {ob vzdol` naprave za uliva-
nje slabov v 13 posameznih podro~jih
cal model of the temperature field2 taking into account a
real caster that has many types of nozzles with various
settings positioned inside a closed cage. A real caster
contains a total of 8 nozzle types and geometrical lay-
outs. The aim is to modify the secondary cooling zones
6, 8, and 10 so as to increase the surface temperature of a
slab in a small radius at the point of the straightening.
Currently, the Lechler 100.638.30.24 air-mist nozzles are
installed in the cooling zones 6, 8, and 10 (Figure 2).
2 MODEL OF THE TEMPERATURE FIELD OF A
SLAB
The presented in-house model of the transient-tem-
perature field of the blank from a slab caster (Figure 1)
is unique as, in addition to being entirely 3D, it can work
in real time. It is possible to adapt its universal code and
apply it to any slab caster. The numerical model covers
the temperature field of the complete length of the blank
(i.e., from the meniscus inside the mould all the way
down to the cutting torch) with up to one million nodes.
The temperature field of the slab passing through a
radial caster with a large radius can be simplified with
the Fourier-Kirchhoff equation, where only the vz com-
ponent of the velocity is considered.



⋅ = ⎛⎝
⎜ ⎞
⎠
⎟ +
⎛
⎝
⎜
⎞
⎠
⎟ + ⎛⎝
⎜ ⎞c
T
x
k
T
x y
k
T
y z
k
T
z
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂ ⎠
⎟ +
+ ⋅ ⋅ +c v
T
z
Qz
∂
∂

source (1)
Equation (1) must cover the temperature field of the
blank in all three stages: above the liquidus temperature
(i.e., the melt), the interval between the liquidus and
solidus temperatures (i.e, the so-called mushy zone) and
beneath the solidus temperature (i.e., the solid phase). It
is therefore convenient to introduce the thermodynamic
function of specific volume enthalpy Hv = cT, which is
dependent on the temperature and also includes the
phase and structural heats (Figure 3).
Heat conductivity k, specific heat capacity c and
density  are thermophysical properties that are also the
functions of temperature. Equation (1) therefore takes
the following form:
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
H
x
k
T
x y
k
T
y z
k
T
z
v


= ⎛⎝
⎜ ⎞
⎠
⎟ +
⎛
⎝
⎜
⎞
⎠
⎟ + ⎛⎝
⎜ ⎞
⎠
⎟ +v
H
zz
v∂
∂
(2)
The unknown enthalpy of the general node of the
blank in the following instant (
 + 
) is given by the
explicit formula:
H H QZ QZ QY QYvi j k vi j k i j i j i j i, ,
( )
, ,
( )
, , , ,(


 
= + + + +1 1 j
QX QX
x y z
+
+ + ⋅
⋅ ⋅
1 )


  
(3)
Figure 3 indicates how the temperature model for the
calculated enthalpy in equation (3) determines the un-
known temperature3. All the thermodynamic properties
of cast steel, dependent on its chemical composition and
the cooling rate, enter the calculation as the functions of
temperature.3 This is therefore a significantly non-linear
task because, even with the boundary conditions, their
dependence on the surface temperature of the blank is
considered here.
The boundary conditions are, therefore, as follows:
1. T T= cast the level of steel (4a)
2. − =k
T
n
∂
∂
0 the plane of symmetry (4b)
3. − = ⋅ −k
T
n
htc T T
∂
∂
( )surface mould inside the mould (4c)
J. STETINA et al.: MINIMIZATION OF SURFACE DEFECTS BY INCREASING THE SURFACE TEMPERATURE ...
312 Materiali in tehnologije / Materials and technology 47 (2013) 3, 311–316
Figure 2: Diagram of the measurement configuration of the cooling
effects of a nozzle
Slika 2: Prikaz konfiguracije meritve hladilnih u~inkov {obe
Figure 3: Enthalpy function for steel showing the phase and structural
changes
Slika 3: Krivulja entalpije jekla, ki prikazuje fazne premene in spre-
membe v mikrostrukturi
4. − = ⋅ − + −k
T
n
htc T T T T
∂
∂
( ) ( )surface amb surface
4
amb
4
within the secondary and tertiary zones (4d)
5. − =k
T
z
q
∂
∂
 beneath the rollers (4e)
The boundary conditions are divided into the area of
the mould, the area of the secondary cooling and the area
of the tertiary cooling.
The initial condition for the investigation is the
setting of the temperature in individual points of the
mesh. A suitable temperature is the highest possible
temperature, i.e., the pouring temperature. The explicit
difference method is used for solving this problem. The
characteristic of this method is that the stability of the
calculation is dependent on the magnitude of the time
step. The model uses a method for adapting the time
step, i.e., the time step entered by an operator is merely a
recommendation and the software is modifying it
throughout the calculation.4
3 HEAT-TRANSFER COEFFICIENT OF THE
NOZZLE
The cooling by the air-mist water nozzles has the
main influence and it is, therefore, necessary to establish
the relevant heat-transfer coefficient of the forced con-
vection. Commercially sold models of the temperature
field describe the heat-transfer coefficient beneath the
nozzles as a function of the incident quantity of water
per unit area. They are based on various empirical
relationships. However, this procedure is undesirable.
The model discussed in this paper obtains its heat-trans-
fer coefficients from the measurements of the spraying
characteristics of all the nozzles used by the caster on the
so-called hot plate in an experimental laboratory and for
a sufficient range of operational pressures of water and a
sufficient range of casting speeds of a slab. This
approach represents a unique combination of an experi-
mental measurement in a laboratory and a numerical
model for the calculation of the non-linear boundary
conditions beneath the cooling nozzle.
A laboratory device enables separate measurements
of individual nozzles. It includes a steel plate mounted
with 18 thermocouples, heated by an external electric
source. The steel plate is heated to the testing tempera-
ture, than it is cooled by a cooling nozzle. On the return
move the nozzle is covered with a deflector, which
enables the movement of the nozzle without cooling the
surface. This device measures the temperatures beneath
the surface of the slab – again by means of thermo-
couples.5
The laboratory device allows the setting of:
• the nozzle type,
• the flow of water,
• the air pressure,
• the distance between the nozzle and the investigated
surface,
• the surface temperature,
• the shift rate.
Since the cooling nozzle 100.638.30.24 for the
minimum water flow appeared to be too intense, the
measurements were made for the smaller nozzle
100.528.30.24 (Figure 4).
Based on the temperatures measured in dependence
of the time, the heat-transfer coefficients (htc) are cal-
culated with an inverse task. They are then processed
further using an expanded numerical and identification
model and converted to the coefficients of function
htc(T,y,z) (Figure 5), which express htc in dependence of
the surface temperature and also the position of the con-
casting with respect to the nozzle. The Lechler air-mist
nozzles show a low dependence of the heat-transfer
coefficient on the slab surface temperature. The value of
htc on the surface of a slab, as it enters the secondary-
cooling zone, significantly affects the process simulation
with respect to the temperature field, the metallurgical
length and also the other technological properties. It,
therefore, affects the prediction of the quality of a slab.
J. STETINA et al.: MINIMIZATION OF SURFACE DEFECTS BY INCREASING THE SURFACE TEMPERATURE ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 311–316 313
Figure 5: Heat-transfer coefficient as a function of the surface tempe-
rature
Slika 5: Koeficient prenosa toplote v odvisnosti od temperature po-
vr{ine
Figure 4: Characteristics of the Lechler nozzles 100.638.30.24 and
100.528.30.24
Slika 4: Zna~ilnosti {obe Lechler 100.638.30.24 in 100.528.30.24
In order to be able to simulate this boundary condi-
tion within the numerical model as accurately as possi-
ble, it is necessary to conduct an experimental measure-
ment on each nozzle in the secondary-cooling zone.
Each of the eight nozzles was measured separately
with the hot model, on which the hot surface of the slab,
cooled by a moving nozzle, can be modelled. The tem-
peratures measured on the surface of the model can be
entered into an inverse task to calculate the intensity of
spraying, which, in turn, can determine htc with a special
mathematical method.
Figure 5 presents the measured values of the heat-
transfer coefficients processed by the temperature model
software. For the nozzle configuration, there is a graph
of the 3D graph (Figure 6) of the heat-transfer coeffi-
cient beneath the nozzle. These graphs are plotted for
surface temperatures from 800 °C to 1000 °C.
4 TEMPERATURE FIELD
The setting of the secondary cooling and its opti-
mization is a very complicated problem. The graph in
Figure 7 shows the resultant temperature fields for indi-
vidual cooling curves. This basic set of graphs serves the
user making it possible to assess which of the cooling
J. STETINA et al.: MINIMIZATION OF SURFACE DEFECTS BY INCREASING THE SURFACE TEMPERATURE ...
314 Materiali in tehnologije / Materials and technology 47 (2013) 3, 311–316
Figure 6: Heat-transfer coefficient for the Lechler air-mist nozzle: a) nozzle 100.638.30.24 with the water flow of 2.2 L/min and air pressure of
0.2 MPa, b) nozzle 100.528.30.24 with the water flow of 2.2 L/min and air pressure of 0.2 MPa, c) nozzle 100.528.30.24 with the water flow of
1.5 L/min and air pressure of 0.2 MPa
Slika 6: Koeficient prenosa toplote pri Lechler {obi za ustvarjanje zra~ne megle: a) {oba 100.638.30.24, pretok vode 2,2 L/min, zra~ni tlak 0,2
MPa, b) {oba 100.528.30.24, pretok vode 2,2 L/min, zra~ni tlak 0,2 MPa, c) {oba 100.528.30.24, pretok vode 1,5 L/min, zra~ni tlak 0,2 MPa
Figure 7: Temperature history along the caster for different configurations of secondary cooling in zones 6, 8 and 10:
a) nozzle 100.638.30.24 with the water flow of 2.2 L/min per nozzle, b) nozzle 100.528.30.24 with the water flow of 2.2 L/min per nozzle, c)
nozzle 100.528.30.24 with the water flow of 1.5 L/min and air pressure of 0.2 MPa
Slika 7: Zgodovina temperature vzdol` livne naprave za razli~no izvedbo sekundarnega hlajenja v conah 6, 8 in 10:
a) {oba 100.638.30.24, pretok vode skozi {obo 2,2 L/min, b) {oba 100.528.30.24, pretok vode skozi {obo 2,2 L/min, c) {oba 100.528.30.24, pretok
vode 1,5 L/min, zra~ni tlak 0,2 MPa
curves is optimal for the given cast steel3. Figure 7a
shows the surface temperature of the slab in the caster
using nozzle 100.638.30.24 with the secondary-cooling
flow of 2.2 L/min per nozzle. Figure 7b shows the tem-
perature for the same conditions, but only for zones 6, 8,
and 10 using nozzle 100.528.30.24. Figure 7c shows the
surface temperature for the water flow of 1.5 L/min for
nozzle 100.528.30.24 in zones 6, 8, and 10. These cal-
culations show that the new nozzles increase the surface
temperature at the straightening point of a small radius
by about 100 °C, while with the higher water-flow rates
a new nozzle can cool as intensely as the original nozzle.
Figure 8 shows the temperature field in the cross-section
at the straightening point using the same parameters as
for Figure 7.
The increase in the surface temperature in the
straightening point on the small-radius surface has
definitely helped to reduce the surface defects of cast
slabs. This conclusion is also confirmed by the macro-
structure figures that were made for two slabs of 1530
mm × 250 mm and a steel grade S275. Figure 9a shows
the macrostructure of the steel that was cast with the
previous setup of the cooling (the Lechler nozzle
100.638.30.24 with the flow rate of 2.2 L/min) and
Figure 9b presents the macrostructure obtained with the
use of the new setup of the cooling (the Lechler nozzle
100.528.30.24 with the flow rate of 1.5 L/min). The new
setup of the cooling has been used on the caster since
July 2012, and therefore the statistical evaluation of the
surface defects from the operational data is not yet
available.
5 CONCLUSIONS
It has been proved that the value of the heat-transfer
coefficient on the surface of the slab and the heat
withdrawal in the secondary-cooling zone significantly
affect the process simulation from the viewpoint of the
temperature field, the metallurgical length and also the
other technological properties. Moreover, these para-
meters also influence the surface quality of cast slabs
and therefore they enable us to predict the quality of the
slabs. In order to simulate the boundary condition of the
numerical model as accurately as possible, it is necessary
J. STETINA et al.: MINIMIZATION OF SURFACE DEFECTS BY INCREASING THE SURFACE TEMPERATURE ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 311–316 315
Figure 9: Macrostructure of the slab before and after the adjustment
of the secondary cooling of steel S 275:
a) nozzle 100.638.30.24 with the water flow of 2.2 L/min per nozzle,
b) nozzle 100.528.30.24 with the water flow of 1.5 L/min per nozzle
Slika 9: Makrostruktura plo{~e pred prilagoditvijo sekundarnega hla-
jenja jekla S 275 in po njej:
a) {oba 100.638.30.24, pretok vode skozi {obo 2,2 L/min,
b) {oba 100.528.30.24, pretok vode skozi {obo 1,5 L/min
Figure 8: Temperature field at the cross-section of the slab at the
straightening point:
a) nozzle 100.638.30.24 with the water flow of 2.2 L/min per nozzle,
b) nozzle 100.528.30.24 with the water flow of 2.2 L/min,
c) nozzle 100.528.30.24 with the water flow of 1.5 L/min
Slika 8: Temperaturno polje na prerezu slaba v to~ki ravnanja:
a) {oba 100.638.30.24, pretok vode skozi {obo 2,2 L/min,
b) {oba 100.528.30.24, pretok vode skozi {obo 2,2 L/min,
c) {oba 100.528.30.24, pretok vode skozi {obo 1,5 L/min
to conduct experimental measurements for each nozzle
in the secondary cooling zone. For this purpose, all the
used nozzles were measured separately with the hot
model, with which the hot surface of the slab, cooled by
a moving nozzle, can be modelled.
The original 3D numerical model of the temperature
field was used for optimizing the surface temperatures of
cast slabs6. The results are presented for the 1530 mm ×
250 mm slabs that are cast in Evraz Vitkovice Steel,
Czech Republic. The performed optimization proved that
a proper replacement of the cooling nozzles in the
secondary cooling zone and a reduction of the flow rate
through the nozzles can help to increase the surface
temperature of a cast slab in the straightening point so
that the surface defects can be reduced. Besides, it was
also shown, for a particular pair of the Lechler nozzles,
that although a nozzle is replaced with a smaller one, in
the cases of higher flow rates, it can have the same
cooling intensity as a larger nozzle.
Acknowledgments
The authors gratefully acknowledge the financial
support from the project GACR P107/11/1566 founded
by the Czech Science Foundation, project ED0002/01/01
of the NETME Centre, and Specific Research FSI-J-12-
22. The co-author, the recipient of the Brno PhD Talent
Financial Aid sponsored by the Brno City Municipality,
also gratefully acknowledges the financial support.
Nomenclature
c specific heat capacity J/kg K
htc heat transfer coefficient W/(m2 K)
Hv volume enthalpy J/m3
k heat conductivity W/(m K)
T temperature K
Tamb ambient temperature K
Tcast melt temperature K
Tsurface temperature in unbending part K
q specific heat flow W/m2
QX, QY, QZ heat flows W
Qsource internal heat source W/m3
x, y, z axes in given direction m
vz casting speed in given direction m/s
 density kg/m3
 Stefan-Bolzmann constant W/m2 K4
 emissivity –

 time s
6 REFERENCES
1 J. P. Birat, The Making, Shaping and Treating of Steel, Casting
Volume, 11th edition, Alan W. Cramb (ed.), The AISE Steeel Foun-
dation, Pittsburgh, PA, USA, 2003
2 T. Mauder, J. Stetina, C. Sandera, F. Kavicka, M. Masarik, An opti-
mal relationship between casting speed and heat transfer coefficients
for continuous casting process, In Metal 2011, Conference proceed-
ings, Metal. Ostrava, Tanger, 2011, 22–27
3 P. Charvat, T. Mauder, M. Ostry, Simulation of latent-heat thermal
storage integrated with room structures, Mater. Tehnol., 46 (2012) 3,
239–242
4 R. Pyszko, M. Prihoda, P. Fojtik, M. Kovac, Determination of heat
flux layout in the mould of continuous casting of steel, Metalurgija
(Metallurgy), 51 (2012) 2, 149–152
5 T. Luks, J. Ondrouskova, J. Horsky, Nozzle cooling of hot surfaces
with various orientations, In Experimental fluid Mechanics 2011,
Proceedings of the International conference, Jicín, 2011, 337–346
6 L. Klimes, P. Popela, An Implementation of Progressive Hedging
Algorithm for Engineering Problems, In Mendel 2010 – 16th Inter-
national Conference on Soft Computing, Brno, BUT, 2010, 459–464
J. STETINA et al.: MINIMIZATION OF SURFACE DEFECTS BY INCREASING THE SURFACE TEMPERATURE ...
316 Materiali in tehnologije / Materials and technology 47 (2013) 3, 311–316
S. M. ALVAREZ et al.: INFLUENCE OF PROCESS PARAMETERS ON THE CORROSION RESISTANCE ...
INFLUENCE OF PROCESS PARAMETERS ON THE
CORROSION RESISTANCE OF CORRUGATED
AUSTENITIC AND DUPLEX STAINLESS STEELS
VPLIV PROCESNIH PARAMETROV NA KOROZIJSKO
ODPORNOST REBRASTIH AVSTENITNIH IN DUPLEKSNIH
NERJAVNIH JEKEL
Sandra M. Alvarez, Asunción Bautista, Francisco Velasco
Universidad Carlos III de Madrid, Materials Sci. and Eng. Dept – IAAB, Avda. Universidad 30, Leganés, Madrid, Spain
asuncion.bautista@uc3m.es
Prejem rokopisa – received: 2012-09-06; sprejem za objavo – accepted for publication: 2012-11-13
The main objective of this work is to study the influence of the forming process on two corrugated, lean, duplex stainless steels
(DSSs): UNS S32001 and UNS S32304. Both grades have been recently proposed as alternative materials to the austenitic UNS
S30403 grade for manufacturing reinforcement bars to be embedded in concrete structures, exposed to corrosive environments.
Hot-worked (HW) corrugated bars of both DSSs are analyzed and their corrosion behaviour is compared with that of the HW
and cold-worked (CW) corrugated bars of S30403.
The corrosion performance is characterized through cyclic polarization curves in 8 different solutions that simulate those
contained inside the pores of concrete in different circumstances.
The obtained results justify a great interest in the studied lean DSS grades with respect to their use as reinforcements. Moreover,
it is proved that the corrugated surface of a bar is clearly less corrosion resistant than the centre of the bar. The processing
method of producing reinforcements influences not only the pitting susceptibility but also the pitting morphology.
Keywords: corrosion, stainless steel, reinforcements, processing, lean duplex
Glavni cilj tega dela je {tudija vpliva postopka izdelave na dve rebrasti dupleksni nerjavni jekli brez molibdena (DSS): UNS
S32001 in UNS S32304. Obe vrsti jekla sta bili predlo`eni kot nadomestni material za avstenitno jeklo UNS S30403 za izdelavo
palic za oja~anje, ki se vgradijo v strukture iz betona, ki so izpostavljene korozivnemu okolju. Vro~e izdelane (HW) rebraste
palice obeh DSS so bile analizirane in njihovo korozijsko vedenje je primerjano z vro~e (HW) in hladno izdelanimi (CW)
rebrastimi palicami iz jekla S30403.
Korozijsko vedenje je ocenjeno z uporabo krivulj pri cikli~ni polarizaciji v 8 razli~nih raztopinah, ki simulirajo teko~ine, ki so v
porah betona v razli~nih okoli{~inah.
Dobljeni rezultati potrjujejo veliko zanimanje za uporabo preu~evanih jekel DSS z nizko vsebnostjo molibdena za oja~itev
betona. [e ve~, dokazano je, da je rebrasta povr{ina palic nedvomno bolj korozijsko odporna kot sredina palice. Metoda izdelave
oja~itvenih palic vpliva poleg ob~utljivosti na jami~asto korozijo tudi na morfologijo korozijskih jamic.
Klju~ne besede: korozija, nerjavno jeklo, oja~itve, izdelava dupleksnega jekla z malo Mo
1 INTRODUCTION
The use of stainless-steel corrugated bars instead of
carbon steel bars in those parts of reinforced concrete
structures that are more exposed to corrosion is one of
the most reliable strategies for assuring the durability of
a structure.1 Initially, austenitic grades were used with
this objective.2
Duplex stainless steels (DSSs) have shown a good
corrosion resistance in many media. With respect to the
reinforced concrete exposed to aggressive environments,
corrosion studies have shown advantages of the tradi-
tional UNS S32205 DSS reinforcements in comparison
with the most common austenitic grades3,4 and they
started to be used for corrugated bars about ten years
ago. Recently, two more economical DSSs have been
proposed for their use in concrete,5 not for replacing the
S32205 reinforcements used in extremely aggressive
conditions, but as an alternative to the corrugated bars of
the austenitic UNS S30403 grade.
UNS S32304 DSS, considered in the present study,
has a low Mo-content and it has been known for years.
Since 2003 its use has been growing in desalinization
industries, marine applications or production processes,
replacing the austenitic UNS S31603 steel. UNS S32001
is the other DSS evaluated in this study. It is a very novel
DSS grade, lower alloyed than S32304, with smaller Ni
and Cr contents, i.e., cheaper.
There are scarce references about the corrosion beha-
viour of these lean DSS grades and most of the studies
consider the environments very different form concrete.
Special chemical characteristics of the solution inside
concrete pores introduce factors that modify the protec-
tive ability of the passive layers on stainless steels. These
characteristics are different from the ones shown by the
materials that are exposed to the atmosphere or to other
environments.6 Most authors also agree that a higher
alkalinity in a pore solution has a positive impact on the
corrosion behaviour of stainless steel,4–8 though the issue
remains controversial.9
Materiali in tehnologije / Materials and technology 47 (2013) 3, 317–321 317
UDK 669.14.018.8:620.19 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)317(2013)
Moreover, to be used as reinforcements in concrete
structures, the stainless steels must be hardened during
the processing10 and their surface must be formed into
corrugations to assure a good adherence with the con-
crete. There are factors that many of the previous
corrosion studies on stainless steels in simulated pore
solutions have not considered, as they were carried out in
stainless steels that were not formed as corrugated
bars.11–14 However, recent studies suggest that the form-
ing process of corrugated bars can dramatically affect the
corrosion behaviour of austenitic stainless steels in
alkaline solutions with chlorides.15
2 EXPERIMENTAL WORK
Four different stainless-steel grades were considered
in the study, two corrugated, lean, duplex stainless steels
(DSSs): the UNS 32001 and UNS S32304 hot-worked
(HW) corrugated bars and two austenitic UNS S30403,
HW and cold-worked (CW) corrugated bars. The pro-
ducts were manufactured by Roldán S. A. (Acerinox
Group, Spain). The diameters of the corrugated bars con-
sidered in the study as well as their chemical composi-
tions can be seen in Table 1. The chemical compositions
of the bars were experimentally determined with X-ray
fluorescence (XRF), using a Spectre XEPOS equipment.
The corrosion behaviour of different places (core and
surface) of the corrugated stainless steels was characte-
rized with cyclic polarizations curves, using an EG&G
263A galvanostat- potentiostat from Princeton Applied
Research. Electrochemical measurements were carried
out in the solutions that simulate those contained in the
concrete pores in different conditions. Saturated
Ca(OH)2 solutions (pH  13), simulating non-carbo-
nated concrete, with four different NaCl contents in mass
fractions were used: (0, 0.5, 1 and 5) %. The saturated
Ca(OH)2 solutions, whose pH values decreased to about
9 due to CO2-bubbling, were used to simulate the
behaviour in carbonated concrete. Chloride contents of
(0, 0.5, 1 and 5) % were also considered for carbonated
solutions.
The testing procedure was based on the ASTM G61
Standard. Cyclic polarization curves were carried out
using a three-electrode cell. A saturated calomel elec-
trode (SCE) was used as the reference electrode and a
stainless-steel mesh as the counter-electrode. Samples of
the corrugated stainless-steel bars acted as working elec-
trodes. The measurements were carried out after a 48-h
exposure of the stainless-steel samples to the testing
solution to assure the correct stabilization of the corro-
sion potential (Ecorr). The sweeping rate was 0.17 mV/s.
The potential was reversed when the current densities
reached a value of 10–4 A/cm2.
To study the corrosion behaviour of a corrugated
surface, samples 2 cm of the real surfaces of the bars
were exposed to the corresponding testing media. The
corrosion behaviour of the non-corrugated materials was
analyzed exposing the samples from the centres of the
bars to the testing solutions. For the samples from the
centres of the bars, an Avesta cell was used to assure the
absence of crevices that could interfere with the measu-
rements.
An analysis of the morphology of the attack after the
polarization curves was carried out with scanning elec-
tronic microscopy (SEM) using a Philips XL30 equip-
ment.
3 RESULTS AND DISCUSSION
The polarization curves of the corrugated surfaces
and the centres of the stainless-steel bars clearly exhibit
different shapes, as can be seen in Figure 1, where the
curves corresponding to the non-carbonated solutions
with 5 % NaCl are shown. For the samples without
corrugations (Figure 1a), the pitting potential (Epit) is
well defined and corresponds with very sharp current
increases. On the other hand, in the tests carried out on
the real surfaces of the bars (Figure 1b), the current
increase after Epit is less pronounced. It must be pointed
out that on certain materials exposed to particular testing
conditions, no corrosion occurred during the test. This is
the case, for example, of the centre of the HW S32304
bar in a non-carbonated solution with 5 % NaCl (Figure
1a), where the current increase does not correspond to
any corrosion phenomenon but it is due to the water
decomposition through reaction 1:
4OH–  2H2O + O2 + 4e
– (1)
In the case of a sudden current increase and the
absence of hysteresis during the reverse cycle, the
potential value confirms that no corrosion has taken
place during the test.
The Epit – Ecorr distance is widely considered to be a
reliable way of measuring the resistance to localized
corrosion. The Ecorr values of all the systems considered
in this study are very similar. The Epit values plotted in
S. M. ALVAREZ et al.: INFLUENCE OF PROCESS PARAMETERS ON THE CORROSION RESISTANCE ...
318 Materiali in tehnologije / Materials and technology 47 (2013) 3, 317–321
Table 1: Diameters of corrugated bars and experimentally determined chemical compositions of the studied stainless steels
Tabela 1: Premer rebrastih palic in eksperimentalno dolo~ena kemijska sestava preiskovanih nerjavnih jekel
Stainless steel Diameterd/mm
Main alloying elements, w/%
S Si Mn Cr Ni Mo N C Fe
CW S30403 10 0.001 0 0.361 1.45 18.30 8.68 0.27 0.050 0.023 Bal.
HW S30403 16 0.001 2 0.298 1.42 18.37 8.74 0.27 0.055 0.026 Bal.
HW S32001 16 0.001 0 0.681 4.14 19.98 1.78 0.24 0.124 0.025 Bal.
HW S32304 16 0.002 0 0.651 1.54 22.70 4.47 0.26 0.153 0.017 Bal.
Figures 2 and 3 can be an adequate tool for comparing
the corrosion behaviours of stainless steels in different
conditions.
It is very interesting to find a significant decrease in
the corrosion resistance of the bars due to the changes to
the corrugations taking place during the forming process.
Figure 2 shows the difference between the Epit values of
the studied stainless steels in carbonated solutions, with
the measurements carried out on the corrugated surfaces
or in the centres of the bars. As in some cases the
definition of Epit is not easy, the potential, at which the
anodic current reaches the value of 10–4 A/cm2, has been
chosen as the criterion for determining this parameter.
The marked difference between the Epit values,
corresponding to the corrugated surface and to the other
regions of a bar, emphasises the effect of the process
parameters. It would be risky to extrapolate the results of
the stainless steels processed in the way different from
that of the corrugated bars to the performance in con-
crete, though it has been often done in literature. These
data confirm the trend observed in the recently published
work on more traditional austenitic stainless steels.15 The
minor corrosion resistance of a corrugated surface of
stainless steel has been explained with a more deformed
microstructure and a higher stress concentration in the
corrugation than found in the centre of the bar.15 Diffe-
rent grain sizes and grain morphologies of the corruga-
tions and of the centres of the bar were studied pre-
viously for the reinforced bars considered in our earlier
work,16 and the obtained results proved that the corru-
gations exhibit a highly deformed microstructure with a
reduced grain size.
In the carbonated solutions without chlorides, no
corrosion was detected for any of the studied stainless
steels. Moreover, the centres of the HW S32304 bars
proved to be immune to corrosion during the polarization
tests carried out independently of the chloride content of
the carbonated solution. However, during the polariza-
tion of the corrugated surfaces of HW S32304 in the
presence of chlorides, current increases corresponding to
a corrosive attack were detected, even with 0.5 % NaCl.
For the other three studied stainless-steel grades, a
localized corrosion always occurred during the polari-
zation tests in the carbonated solutions with chlorides.
S. M. ALVAREZ et al.: INFLUENCE OF PROCESS PARAMETERS ON THE CORROSION RESISTANCE ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 317–321 319
Figure 3: Differences between the Epit measured for the centres of
stainless-steel bars and for the corrugated surfaces in non-carbonated
Ca(OH)2 solutions with different chloride contents. Conditions with-
out the plotted Epit values correspond to tests where no corrosion takes
place.
Slika 3: Razlika med izmerjeno Epit za sredino nerjavne palice in za
rebrasto povr{ino v negazirani raztopini Ca(OH)2 z razli~no vsebnost-
jo kloridov. Primeri brez prikazane Epit ustrezajo preizkusom, kjer se
korozija ni pojavila.
Figure 1: Polarization curves in non-carbonated Ca(OH)2 solutions
with 5 % NaCl: a) centres of the bars, b) corrugated surfaces
Slika 1: Polarizacijske krivulje v negazirani raztopini Ca(OH)2 s 5 %
NaCl: a) sredina palice, b) rebrasta povr{ina
Figure 2: Differences between the Epit measured in the centres of
stainless-steel bars and on corrugated surfaces in carbonated Ca(OH)2
solutions with different chloride contents
Slika 2: Razlika med izmerjenim Epit v sredini palice iz nerjavnega
jekla in na rebrasti povr{ini v gazirani raztopini Ca(OH)2 z razli~no
vsebnostjo kloridov
For all the materials the Epit of the corrugated surface is
much lower than the Epit of the centre of a bar.
As expected, in all the cases an increase in the
chloride content of the solution causes a decrease in the
resistance to localized corrosion, i.e., a decrease in Epit.
In Figure 3, the Epit values detected in non-carbo-
nated solutions are plotted. It can be seen that, at a
higher pH, it is more difficult to cause corrosion during
the test. In the solutions without chlorides, no corrosion
occurs in any of the cases and the centre of HW S32304
is immune to the attacks in the testing media with
chlorides, as reported for pH  9. Besides, no corrosion
occurs during the polarization in the solutions with 0.5 %
NaCl on the centres of the other studied bars. In the case
of the 1 % NaCl testing solution, no corrosion was found
on the centres of HW S30403 or HW S32001.
The corrugated surfaces of the bars prove again to be
much more prone to corrosion. For this type of samples,
the only condition where no pitting is detected is HW
S32304 with 0.5 % NaCl. The important influence of the
microstructural changes occurring in the surfaces of the
corrugated bars during the forming process is again
clearly proved.
If the results from Figures 2 and 3 are used to
compare the corrosion behaviours of different grades, it
is demonstrated that HW S32304 is clearly more
corrosion resistant than any of the studied austenitic
grades. Despite the volatility of the prices in the market,
it can be considered that a S32304 grade can cost about 9
% less than a S30403 grade. This result justifies the great
interest in this DSS grade that is seen as an alternative
for the traditional austenitic grade used in these
applications, as S32304 has a better performance and it
is somewhat more economical. The DSS S32001 grade
can be estimated to be about 15 % cheaper than the
austenitic S30403. The results of the corrosion tests
carried out indicate that the corrosion resistance of both
grades are quite similar, or that the corrosion resistance
of the cheap, new DSS grade is even better.
In addition to Epit, another interesting parameter,
which can be obtained from the polarization curves, is
the maximum intensity (imax) reached during the
measurements. All the curves are programmed to reverse
the potential sweep when a current intensity of 10–4
A/cm2 is reached. When no corrosion occurs, imax is 10–4
A/cm2, as the current quickly decreases when the applied
potentials decrease. When pits are formed during the
anodic polarization, the current still increases as the
potentials start to decrease due to the important
autocatalytic effect of the localized corrosion. The higher
the imax, the more dangerous is the pitting morphology.
As an example, the values of this parameter for sixteen
tested conditions of HW S32001 were plotted in Figure
4. If the conditions, under which no corrosion takes
S. M. ALVAREZ et al.: INFLUENCE OF PROCESS PARAMETERS ON THE CORROSION RESISTANCE ...
320 Materiali in tehnologije / Materials and technology 47 (2013) 3, 317–321
Figure 5: Images of different morphologies of the attacks that appear
on the centre of a bar and on the corrugated surface. The pits after the
polarization of HW S32001 in non-carbonated Ca(OH)2 solutions with
5 % NaCl: a) corrugated surface, b) center of the bar.
Slika 5: Posnetki razli~nih morfologij napada, ki se pojavi v sredini
palice in na rebrasti povr{ini. Jamice po polarizaciji HW S32001 v
negazirani raztopini Ca(OH)2 s 5 % NaCl: a) rebrasta povr{ina, b) sre-
dina palice.
Figure 4: Differences between the imax values obtained with the
polarization curves for the centres and the corrugated surfaces of
duplex HW S32001 bars in carbonated and non-carbonated Ca(OH)2
solutions with different chloride contents
Slika 4: Razlike med vrednostmi imax, dobljene iz polarizacijskih
krivulj iz sredine in iz rebraste povr{ine dupleksnih palic HW S32001
v gazirani in negazirani raztopini Ca(OH)2 z razli~no vsebnostjo
kloridov
place, are not considered (0 % NaCl and the centre of the
bar for 0.5 and 1 % NaCl at pH  13), it can be seen that
the samples from the centre of the bar, though less
susceptible to corrosion than the corrugated surfaces,
suffer from a more aggressive attack than when it occurs
on the surface. The same conclusion is reached if the
results for the other four studied materials are analyzed.
It can also be seen in Figure 4 that chlorides have an
important influence on the increase of imax.
An observation of the morphology of the pits after
the polarization curves confirms the idea deduced from
the imax values. As it can be seen in Figure 5, the
polarization causes small, shallow pits widely distributed
on the most deformed regions of the surface of the
corrugation. In the centre of the bar, polarizations cause
scarce, but much bigger pits that can be much more
dangerous.
4 CONCLUSIONS
The susceptibility to pitting corrosion on the corru-
gated surface of corrugated stainless steel is always
much higher than in the centre of the bars of the same
material. The forming process clearly decreases the
corrosion resistance of stainless steel used as a reinforce-
ment material in concrete structures.
The attack that appears on the corrugated surfaces of
stainless steels during an anodic polarization is less
localized and less dangerous than the attack that appears
in the centres of the bars.
The new lean DSSs for reinforcing bars are very
interesting options for substituting the traditional auste-
nitic S30403 bars. S32304 clearly exhibits a better corro-
sion behaviour being also somewhat cheaper. S32001 is
highly interesting from an economic point of view and
its corrosion results are similar, even slightly better, than
those of S30403.
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15 E. C. Paredes, A. Bautista, S. M. Alvarez, F. Velasco, Corrosion
Science, 58 (2012), 52–61
16 S. M. Alvarez, A. Bautista, F. Velasco, S. Guzman, Proceedings of
the 7th European Stainless Steel Conference, Como, Italy, 2011
S. M. ALVAREZ et al.: INFLUENCE OF PROCESS PARAMETERS ON THE CORROSION RESISTANCE ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 317–321 321

S. KÜÇÜKDERMENCI et al.: CONCENTRATION AND PATH-LENGTH DEPENDENCE ON THE FARADAY ROTATION ...
CONCENTRATION AND PATH-LENGTH DEPENDENCE
ON THE FARADAY ROTATION OF MAGNETIC FLUIDS
BASED ON HIGHLY WATER-SOLUBLE Fe3O4/PAA
NANOPARTICLES SYNTHESIZED BY A
HIGH-TEMPERATURE HYDROLYSIS METHOD
ODVISNOST KONCENTRACIJE IN DOL@INE POTI OD
FARADAYEVEGA VRTENJA MAGNETNIH TEKO^IN NA OSNOVI
VISOKOVODOTOPNIH NANODELCEV Fe3O4/PAA,
SINTETIZIRANIH Z METODO VISOKOTEMPERATURNE
HIDROLIZE
Serhat Küçükdermenci1,5, Deniz Kutluay2, Erdal Çelik3,4, Ömer Mermer1,
Ýbrahim Avgýn1
1Department of Electrical and Electronics Engineering, Ege University, Bornova 35100, Izmir, Turkey
2Department of Electronics and Communications Engineering, Izmir University, Uckuyular 35290, Izmir, Turkey
3Center for Fabrication and Applications of Electronic Materials (EMUM), Dokuz Eylul University, Týnaztepe Campus, 35160, Izmir, Turkey
4Department of Metallurgical and Materials Engineering, Dokuz Eylul University, Týnaztepe Campus, 35160, Izmir, Turkey
5Department of Electrical and Electronics Engineering, Balikesir University, Campus of Cagis 10145, Balikesir, Turkey
serhat.kucukdermenci@ege.edu.tr
Prejem rokopisa – received: 2012-09-13; sprejem za objavo – accepted for publication: 2012-11-16
In this study, highly water-soluble Fe3O4/PAA (polyacrylic acid) nanoparticles (NPs) were synthesized by a high-temperature
hydrolysis method. We report the first demonstration of the concentration and path-length dependence on the Faraday rotation
(FR) for a magnetic fluid (MF) synthesized by this novel method. Experiments were performed in the DC regime (0–6 × 10–2 T)
at room temperature. Measurements were carried out with 5–3.33 mg/ml and 1.18 mg/ml samples in cells (2, 5, 7 and 10) mm.
The maximum rotation was recorded as 0.96° cm–1 for the 3.33 mg/ml concentration in the cell 10 mm. It was found that the
magnetic fluid behaves with a distinctive phenomenon in different sized cells although its concentration was the same. The role
of the different parameters on the FR was discussed via spatial limitations imposed by the cells and a premature saturation term.
This work provides a new insight for FR investigations of MFs including highly water-soluble magnetic NPs.
Keywords: magnetic fluids, water soluble Fe3O4/PAA-nanoparticles, Faraday rotation
V tej {tudiji so bili sintetizirani visokovodotopni Fe3O4/PAA (poliakrilna kislina)-nanodelci (NPs) z metodo visokotemperaturne
hidrolize. Poro~amo o prvih predstavitvah odvisnosti koncentracije in dol`ine poti od Faradayevega vrtenja (FR) za magnetno
teko~ino (MF), sintetizirano po tej novi metodi. Preizkusi so bili izvedeni v DC-re`imu (0–6 × 10–2 T) pri sobni temperaturi.
Meritve so bile izvr{ene z vzorci 5–3,33 mg/ml in 1,18 mg/ml v celicah (2, 5, 7 in 10) mm. Maksimum vrtenja je bil ugotovljen
kot 0,96° cm–1 pri koncentraciji 3,33 mg/ml v celici 10 mm. Ugotovljeno je bilo, da magnetna teko~ina izkazuje zna~ilne pojave
v celicah razli~ne velikosti, ~eprav je bila njena koncentracija enaka. Vloga razli~nih parametrov na FR je bila razlo`ena s
prostorskimi omejitvami, ki jo povzro~ijo celice in z uvajanjem izraza za prezgodnjo prenasi~enost. To delo zagotavlja nov
vpogled v preiskave Faradayevega vrtenja (FR) magnetnih teko~in (MFs), vklju~no z visokovodotopnimi magnetnimi nanodelci
(NPs).
Klju~ne besede: magnetne teko~ine, vodotopni Fe3O4/PAA-nanodelci, Faradayevo vrtenje
1 INTRODUCTION
Magneto-optic (MO) effects occur in gases, liquids,
and solids. In general, solids exhibit the strongest MO
effects, liquids exhibit weaker effects, and gases exhibit
the weakest effects.1 FR in MFs has been demonstrated
in the visible,2 near-infrared,3,4 and mid-infrared5 regi-
mes. The wavelength dependence of FR and Faraday
ellipticity was investigated by N. A. Yusuf and co-wor-
kers.6 MO experiments were performed on size-sorted
iron oxide (-Fe2O3) MFs.7 Water-based MF samples
were synthesized by the coprecipitation method followed
by a size-sorting process. Davies and Llewellyn8 have
reported measurements on FR in highly diluted samples
of Fe3O4 and CO particle MFs. From their results they
concluded that FR is governed by magnetization. Yusuf
et al.9 have reported measurements on FR in a relatively
concentrated sample.
The MO response of MFs depends on various para-
meters including particle size, the concentration of the
particles, the externall applied magnetic field, tempera-
ture, and the kinds of particles, surfactants, and carrier
fluids.10 Due to the large variety of MFs, measurements
of the MO effects of MFs have still not been reported so
widely for novel ones and each new sample should be
carefully analyzed.
In 2007, the Yin group synthesized novel superpara-
magnetic magnetite colloidal NPs which can self-assem-
Materiali in tehnologije / Materials and technology 47 (2013) 3, 323–327 323
UDK 66.017:537.622 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)323(2013)
ble into one-dimensional (1D) particle chains and exhibit
excellent tunable photonic properties.11,12 Since then,
there has been a widespread interest in these NPs and
their applications. Despite their tremendous potential in
various applications, interesting fundamental questions
which refer to their colloidal crystallization with and
without a magnetic field remain unanswered. Therefore,
in this paper we report the first demonstration of the con-
centration and path-length dependence on FR for MF
based on these NPs. Here we report the results of an
investigation of long-term stable MFs including highly
water-soluble NPs showing FR in the DC regime (0–6 ×
10–2 T) at room temperature. Water-based MF samples
were synthesized by a novel high-temperature hydrolysis
method.
The different relationships between the chain length
and the concentration obtained by different workers may
be attributed to the dependence of the chain length on
other parameters such as the field, temperature and spa-
tial limitations imposed by the cells used in containing
the sample under investigation.13–15 In our previous
work16 we measured the FR for 14 different concentra-
tions from 1.8 mg/ml to 5 mg/ml in a cell 10 mm. It was
observed that the FR increases for higher concentrations
(from 1.18 mg/ml to 3.33 mg/ml). The maximum rota-
tion was recorded as 0.96 °/cm for 3.33 mg/ml and this
was named the critical concentration (CCRITICAL). It was
found that rotation tends to decrease when the concentra-
tion is higher than CCRITICAL (for 4 mg/ml and 5 mg/ml).
We chose three samples and repeated the measurements
for 2–5 mm and 7 mm cells additionally for observing
the path-length dependence on the FR and investigating
the spatial limitation factor on MO effects.
2 EXPERIMENTAL DETAILS
The synthesis of Fe3O4 NPs was explained in detail in
our previous report16 and is summarized here. For the
synthesis, diethylene glycol (DEG, 99.9 %), anhydrous
ferric chloride (FeCl3, 97 %), sodium hydroxide (NaOH,
96%), and polyacrylic acid (PAA, Mw = 1800) were ob-
tained from Sigma-Aldrich. Distilled water was used in
all the experiments. All the chemicals were used as
received without an further treatment and/or purification.
For the synthesis of the Fe3O4 NPs a NaOH/DEG solu-
tion was prepared by dissolving 100 mmol of NaOH in
40 m of DEG at 120 °C under nitrogen for 1 h. Then, the
light-yellow solution was cooled to 70 °C (stock solution
A). Using a 100 ml three-necked flask equipped with a
nitrogen inlet, a stirrer and a condenser, 10 mmol of
FeCl3 and 20 mmol of PAA were dissolved in 41 ml of
DEG under vigorous stirring. In this method, DEG is not
only a reducing agent but also a solvent in the reaction.
The solution was purged with bubbling nitrogen for 1 h
and then heated to 220 °C for 50 min (stock solution B).
Subsequently, 20 ml of NaOH/DEG solution was
injected rapidly into the above solution. The reaction was
allowed to proceed for 2 h. The black color of the solu-
tion confirmed the formation of magnetite nanoclusters.
The resultant black product was repeadedly washed with
a mixture of ethanol and water and collected with the
help of a magnet. The cycle of washing and magnetic
separation was performed five times. A one-pot synthesis
was conducted for the Fe3O4/PAA nanoparticles, and
thus no extra separate process was needed for the surface
modification. The polyol method is based on the theory
that NPs will be yielded upon heating precursors in a
high-boiling-point alcohol at elevated temperature. DEG
is chosen here as the solvent inasmuch as it can easily
dissolve a variety of polar inorganic materials due to its
high permittivity ( = 32) and high boiling point
(246 °C). PAA is used as the capping agent, on which the
carboxylate groups show strong coordination with Fe3+
on the Fe3O4 surface and the uncoordinated carboxylate
groups extend into the water solution, rendering particles
with a high water dispersibility.
Powder X-ray diffraction (XRD) analysis of the
Fe3O4 NPs was performed on a Phillips EXPERT 1830
diffractometer with Cu K
 radiation. The XRD data were
collected over the range of 10–80° (2) with a step inter-
val of 0.02° and a preset time of 1.6 s per step at room
temperature. Particle size distributions of the NPs were
measured using a Zetasizer 4 Nano S dynamic light
scattering (DSL, Malvern, Worcestershire, UK). The
light-scattering measurements were carried out with a
laser of wavelength 633 nm at a 90° scattering angle.
The magnetic measurements were carried out using a
LakeShore 7400 (Lakeshore Cryotronics) vibrating-sam-
ple magnetometer (VSM) at 300 K. For the FR experi-
ments we used a Thorlabs model HGR20 2.0 mW,
543 nm laser source, a GMW Electromagnet Systems
model 5403 electromagnet, a Kepco power supply model
BOP 20-5M, a LakeShore Model 455 DSP Gaussmeter,
a Stanford Research Systems Model SR830 DSP lock-in
amplifier, a New focus model 2051 photo detector, a ILX
Lightwave model OMM – 6810B optical multimeter and
a model OMH – 6703B silicon power head.
3 RESULTS
Figure 1 shows the XRD pattern of PAA-coated
Fe3O4 NPs synthesized by the high-temperature hydroly-
sis method. Note that the (220), (311), (400), (422),
(511), and (440) diffraction peaks observed for the
curves can be indexed to the cubic spinel structure, and
all the peaks were in good agreement with the Fe3O4
phase (JCPDS card 19-0629). It is clear from Figure 2
that the average diameter of the Fe3O4/PAA NPs
obtained from the DLS analysis is approximately 10 nm.
The magnetic properties of the NPs were measured at
300 K using the VSM. As given in Figure 3, the satu-
ration magnetization was determined as 38.8 emu/g.
Worth nothing here is that the NPs showed no remanence
S. KÜÇÜKDERMENCI et al.: CONCENTRATION AND PATH-LENGTH DEPENDENCE ON THE FARADAY ROTATION ...
324 Materiali in tehnologije / Materials and technology 47 (2013) 3, 323–327
or coercivity at 300 K, i.e., they exhibited superparamag-
netic behavior.
The particle content is 20 mg for various concen-
trations with the addition of distilled water from 4 ml to
17 ml. The NPs kept dispersing well after the MF had
been standing for more than 4 weeks and no sedimenta-
tion was observed for all the samples. Maintaining long-
term stability can make these kinds of MFs ideal
candidates for optical devices.
The FR measurements for 14 different concentrations
from 1.8 mg/ml to 5 mg/ml in a 10 mm cell were shown
graphically in Figure 4. The maximum FR of these
measurements was shown in Figure 5. We chose three
concentrations for investigating the effect of cell size on
the FR measurements. The chosen concentrations were:
3.33 mg/ml CCRITICAL, for which we measured the maxi-
mum FR; 1.18 mg/ml, for which we measured the
weakest FR before CCRITICAL; and 5 mg/ml, for which we
measured the lowest FR after CCRITICAL. The measure-
ments were repeated for 2–5 mm and 7 mm cells for
these concentrations.
For sample 14 (Table 1) the FR was found to be less
than 0.1 degree in the 2 mm thick cell. For the samples
that have same concentration but in different cell sizes,
the magnetization must be the same at a particular field
since all have an equal magnetic dipole moment per unit
volume. Therefore, if the FR is only governed by magne-
tization, it would be the same value for these samples. In
order to explain the results, we can introduce terms to the
S. KÜÇÜKDERMENCI et al.: CONCENTRATION AND PATH-LENGTH DEPENDENCE ON THE FARADAY ROTATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 323–327 325
Figure 3: Hysteresis loops of superparamagnetic nanoparticles at
room temperature
Slika 3: Histerezna zanka super paramagnetnih nanodelcev pri sobni
temperaturi
Figure 2: Particle size distribution of Fe3O4 NPs synthesized by
high-temperature hydrolysis
Slika 2: Razporeditev velikosti nanodelcev Fe3O4, sintetiziranih z me-
todo visokotemperaturne hidrolize
Figure 1: XRD pattern of PAA-coated Fe3O4 NPs synthesized by
high-temperature hydrolysis
Slika 1: Rentgenska difrakcija (XRD) s PAA, prekritega z Fe3O4
nanodelci, sintetiziranimi z metodo visokotemperaturne hidrolize Figure 4: FR concentration dependence for 10-mm cell
Slika 4: Odvisnost FR od koncentracije pri celici 10 mm
Figure 5: Maximum FR graph of samples for 10-mm cell
Slika 5: Predstavitev FR maksimumov vzorcev pri celici 10 mm
Faraday rotation which is due to the chain formation and
is given by:
C ( ) ( )B VBl B= (1)
where C(B) is the Faraday rotation due to the chain
formation, B is the local field, l(B) is the chain length at
field B and V is the Verdet constant.
For a homogeneous colloidal system of single-do-
main fine ferromagnetic particles, the Faraday rotation is
given by:
M
S
( )
( )
B C
M B
M
= (2)
in which M(B) is the Faraday rotation at magnetic field
B, M(B) is the magnetization of the sample at magnetic
field B, MS is the saturation magnetization of the sample
and C is a constant. Both the magnetization of the
sample and the chain formation FR for the MFs can be
expressed as shown below:
  ( )
( )
( ) ( ) ( )B C
M B
M
VBl B B B= + = +
S
M C (3)
Here, C is a constant that can be found the at high
field assuming that saturated chain length has no
change.14 This artificial saturation can be referred to as
"premature saturation" and utilized to explain the diffe-
rent behavior of the same sample but in a different cell
thicknesses.
Since samples 1 and 14 were less imposed effect of
chain formation, their initial slopes in various cell thick-
ness do not diverge too much in the 0–2 × 10–2 T. Never-
theless, at the optimum concentration, sample 3 which
has the maximum effect of chain formation, despite the
same amount of NPs, the initial slope in different thick-
ness cells is quite different.
We also observed that the FR of the same concentra-
tion fluids reached saturation at higher fields for a longer
path length (Figure 6). This effect might be the result of
premature saturation phenomena of the mean chain
length in fluids. Notably, for samples 3 and 14, changing
S. KÜÇÜKDERMENCI et al.: CONCENTRATION AND PATH-LENGTH DEPENDENCE ON THE FARADAY ROTATION ...
326 Materiali in tehnologije / Materials and technology 47 (2013) 3, 323–327
Figure 6: FR of three different concentrated MF in different size cells
including: a) sample 1 (concentration: 20 mg/(4 ml)), b) sample 3
(concentration: 20 mg/(6 ml)) and c) sample 14 (concentration: 20
mg/(17 ml)). The dimensions of the cells are indicated on the figures.
Slika 6: FR treh magnetnih raztopin (MF) z razli~no koncentracijo pri
celicah z razli~no velikostjo: a) vzorec 1 (koncentracija: 20 mg/(4
ml)), b) vzorec 3 (koncentracija: 20 mg/(6 ml)) in c) vzorec 14 (kon-
centracija: 20 mg/(17 ml)). Velikosti celic so ozna~ene na slikah.
Table 1: Maximum FR of three different concentrated MF in four
different cells
Tabela 1: Maksimum za FR pri treh razli~nih koncentracijah MF v 4
razli~nih celicah
Sample
number
Concentration
(mg/ml)
FR maximum degree (°/cm)
Cell size (mm)
2 5 7 10
1 5.00 0.3 0.36 0.43 0.49
2 4.00 0.87
3 3.33(CCRITICAL)
0.15 0.37 0.6 0.96
4 2.86 0.83
5 2.50 0.78
6 2.22 0.70
7 2.00 0.65
8 1.82 0.56
9 1.67 0.48
10 1.54 0.44
11 1.43 0.37
12 1.33 0.35
13 1.25 0.30
14 1.18 <0.1 0.13 0.22 0.28
of saturation values on FR with different cell size
indicates premature saturation of the chain length.
4 DISCUSSION
The dipole-dipole interactions can be described as
F = 3μ2 (1 – 3 cos2 
)/l4, where μ is the induced magnetic
moment, 
 is the angle between the dipole and the line
connecting the dipoles, and l is the center-center distance
between two particles. When two dipoles are aligned
head-to-end, the dipole interaction is an attraction Fma =
–6(μ2/l4). While they are aligned side by side, the inte-
raction becomes a net repulsion Fmr = 3(μ2/l4). Along the
magnetic field, the particles attract each other and form
chains due to the head-to-end alignment of dipoles.17
The formed chain length varies with applied mag-
netic field and the concentration of the MF. The FR is
not only governed by the magnetization of the fluid but
also by the chain formation. The dimension of the cell in
the field direction plays an important role in the length
and number of chains formed in the sample. For short
path length, a "spatial limitation" on the chain length is
imposed, resulting in a "premature saturation" of the
length. Jones and Niedoba18 had a drop of undiluted fluid
sandwiched between two parallel glass cover slips and
placed normal to the optical axis of a microscope. The
magnetic field was applied parallel to the axis of the
microscope, i.e., perpendicular to the plane of the
thin-film sample. In this experimental arrangement Jones
and Niedoba observed the number of chains per unit
volume, n, and found it increased rapidly with the
applied field. The experimental set up used by Jones and
Niedoba put a severe limitation on the length of the
formed chains and led to premature saturation in the
chain lengths, as was suggested by Fang et al.19 It is
worth mentioning that for thin samples, the chain length
approaches the thickness of the sample at very low fields
and saturates prematurely. Consequently, the chain
length is not expected to increase any further with the
field but rather an increase in the number of chains is
observed.
If the Faraday rotation is only governed by the
magnetization of the sample, then the saturation will take
place at practically the same applied field for the same
sample, regardless of the thickness of the sample. The
difference in the saturation field is attributed to the chain
formation in the sample. As was mentioned above, this
chain formation may prematurely saturate in thin sam-
ples due to the physical limitation imposed on the
sample.
5 CONCLUSIONS
In summary, high-quality Fe3O4/PAA NPs for MF
formation were successfully synthesized. We report the
first demonstration of the concentration and path-length
dependence on the FR for MF synthesized using this
novel method. We have demonstrated the FR of highly
water-soluble MF measured in the 0–6 × 10–2 in the DC
regime. The effects of both spatial limitations imposed
by the cells and the chain concentration were observed
on the FR. We observed that the chain length is also
determined by the physical dimensions of the sample.
Four different cells were used for the path-length experi-
ments. The FR of the same concentration fluids with
shorter path lengths reached saturation at lower fields as
a result of the premature saturation of an average chain
length. The experimental results shed some light on the
role of agglomeration and chain formation in the FR.
Magneto-optical effects in the magnetic fluids have a
promising potential for technological and industrial
applications as well as their academic importance. These
measurements can help to understand the MO behavior
of the MFs, including highly water-soluble magnetic
NPs.
Acknowledgements
The authors specially wish to thank Dr. Yavuz Öztürk
for many helpful discussions.
6 REFERENCES
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V. Part 3, no. CRC Press, Boca Raton 1995
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(1988), 64
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Int. Ed., 23 (2007) 46, 4342–4345
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IEEE Trans. Magn., 26 (1990), 2852
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71–78
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(2007), 77
S. KÜÇÜKDERMENCI et al.: CONCENTRATION AND PATH-LENGTH DEPENDENCE ON THE FARADAY ROTATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 323–327 327

\. VA[TAG et al.: THE POSSIBILITY OF COPPER CORROSION PROTECTION IN ACIDIC MEDIA ...
THE POSSIBILITY OF COPPER CORROSION
PROTECTION IN ACIDIC MEDIA USING A THIAZOLE
DERIVATIVE
MO@NOST PROTIKOROZIJSKE ZA[^ITE BAKRA V KISLEM
MEDIJU Z UPORABO DERIVATOV TIAZOLA
\en|i Va{tag1, Suzana Apostolov1, Miodrag Had`istevi}2, Milenko Sekuli}2
1University of Novi Sad, Faculty of Sciences, Trg D. Obradovi}a 3, 21000 Novi Sad, Serbia
2University of Novi Sad, Faculty of Technical Sciences, Trg D. Obradovi}a 6, 21000 Novi Sad, Serbia
miodrags@uns.ac.rs
Prejem rokopisa – received:2012-09-14; sprejem za objavo – accepted for publication: 2012-11-22
Assessing financial losses on an annual scale, it was determined that the world economy loses about 2.2 trillion US dollars as a
result of the corrosion of materials. The most significant losses are caused by the corrosion of metals. Against corrosion, metals
can be protected in several ways: by cathodic protection or anodic protection, by the use of inhibitors, or by using various
protective coatings that can be metallic, non-metallic, organic or inorganic. In this paper, the possibility of the corrosion
protection of copper was investigated in acidic media using inhibitors. The inhibitor properties of 5-(4’-dimetylamino-
benzylidene)-2,4-dioxotetrahydro-1,3-thiazole, (DABDT) were tested on copper corrosion in an acidic sulphate-containing
solution (0.1 mol dm–3 Na2SO4, pH = 3). Using potentiostatic polarization measurements, the inhibitor efficiency of DABDT as
a function of concentration was determined and the mechanism of its adsorption on the copper surface was defined. It was found
that the DABDT thiazole derivative acts as mixed inhibitor on copper corrosion in acidic media. In the investigation range,
increasing the concentration results in a better inhibition efficiency of the DABDT.
Keywords: copper, corrosion, thiazole derivatives, polarization measurements
Ocena finan~nih izgub na letnem nivoju je pokazala, da svetovna ekonomija izgubi okoli 2,2 milijardi ameri{kih dolarjev kot
posledico korozije materialov. Najpomembnej{e izgube so povzro~ene s korozijo kovin. Proti koroziji lahko kovine za{~itimo na
ve~ na~inov: s katodno oziroma anodno za{~ito, z uporabo inhibitorjev ali razli~nih za{~itnih prevlek, ki so lahko kovinske,
nekovinske, organskega oziroma anorganskega izvora. V tem prispevku je bila raziskana mo`nost protikorozijske za{~ite bakra
v kislem mediju z uporabo inhibitorjev. Lastnosti inhibitorja iz 5-(4’-dimetilaminobenziliden)-2,4-dioksotetrahidro-1,3-tiazol,
(DABDT), so bile preizku{ene pri koroziji bakra v kislem mediju, ki je vseboval raztopino (0,1 mol dm–3 Na2SO4, pH = 3). S
potenciostati~nimi meritvami polarizacije je bila dolo~ena u~inkovitost DABDT kot funkcija koncentracije in mehanizem
njegove adsorpcije na povr{ino bakra. Odkrili smo, da se DABDT-derivati tiazola vedejo kot me{ani inhibitorji pri koroziji
bakra v kislem mediju. V raziskanem obmo~ju se je z nara{~anjem koncentracije DABDT izbolj{ala u~inkovitost zadr`evanja
korozije.
Klju~ne besede: baker, korozija, derivati tiazola, meritve polarizacije
1 INTRODUCTION
In accordance with recent estimates it has been
calculated that the effect of corrosion on the US
economy in 2012 exceeds 1 trillion dollars per year for
the first time.1 Taking into account this fact, it is clear
that corrosion is the cause of significant losses in the
economy of each country and therefore falls within one
of the important factors of the global financial and
energy crisis. It is understandable, therefore, that the
great interest and tendency to reduce losses as a result of
the corrosion of construction materials is to be kept to a
minimum.
Copper and its alloys have good characteristics and a
wide range of industrial applications. They have the
largest number of applications as conductors of elec-
tricity and heat. Copper shows excellent performance as
a structural material, because among other things, it is
resistant to corrosion over a wide range of pH values.
However, it is known that in aggressive media, copper is
susceptible to corrosion, due to the lack of a protective
passive layer in the acidic environment2. Copper as a
structural material is often exposed to acidic conditions,
e.g., during its purification, electro polishing or during
the removal of corrosion products from the heat trans-
missions. The most commonly used acid in these pro-
cesses is sulphuric acid. The protection of copper in
these processes in an acidic environment is usually
accomplished with corrosion inhibitors.3–5
Organic compounds that contain hetero-atoms such
as nitrogen, sulphur and oxygen, or conjugated double
bonds have shown good inhibition properties against
copper corrosion.6–8 This kind of organic molecules can
be adsorbed at the metal-solution interface, which will
reduce the corrosive attack on the metal in acidic
media.9,10
The degree of corrosion protection of these molecu-
les depends on the strength of the interaction between
the organic molecule and the metal surface atoms.
The thiazole derivatives are an interesting group of
nitrogen- and sulphur-containing organic compounds
Materiali in tehnologije / Materials and technology 47 (2013) 3, 329–333 329
UDK 620.197.3:669.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)329(2013)
that act as inhibitors against the dissolution of copper in
acidic electrolytes.
The aim of this study is to investigate and interpret
how the thiazole derivative 5-(4’ -dimetylaminobenzyl-
idene) -2,4- dioxotetrahydro-1,3-thiazole, (DABDT),
inhibits the copper corrosion in aerated 0.1-M Na2SO4 at
pH = 3.
2 EXPERIMENTAL
Reagent-grade chemicals (Merck) and double-
distilled water were used to prepare the electrolyte of
0.1-M Na2SO4 at pH = 3; the pH was adjusted to 3.0
using diluted sulphuric acid.
The investigated thiazole derivative was 5-(4’-dime-
tylaminobenzylidene)-2,4- dioxotetrahydro-1,3-thiazole
(DABDT) (Figure 1).
Figure 1: Structure of the investigated thiazoles
Slika 1: Struktura raziskanih tiazolov
All the experiments were conducted in the open
atmosphere and at room temperature. Due to the low
solubility of the tested thiazole derivative, the inhibitor
was first dissolved in 20 ml of ethanol.
Weight-loss measurements were carried out on
copper coupons (6 cm × 1.5 cm × 0.2 cm) in 0.1 mol
dm–3 Na2SO4 at pH = 3 for different inhibitor concentra-
tions (0.001 mmol dm–3–0.01 mmol dm–3). The coupons
immersed in the test solutions were allowed to stand for
one week in an air atmosphere.
For the electrochemical measurements a three-elec-
trode cell was used. High-purity copper rods (99.99 %
Cu) with an exposed area of 0.7 cm2 were used as the
working electrode, a saturated calomel electrode (SCE)
was used as the reference electrode and platinum was
used as the counter electrode. Before using, the working
electrode was wet-polished with SiC papers (grit sizes of
800 and 1200), rinsed with acetone and double-distilled
water. All the measurements were made at room tempe-
rature.
The polarization measurements were performed at
five different inhibitor concentrations in the concen-
tration range 0.001–0.01 mmol dm–3. The measurements
were carried out when the open-circuit potential (OCP)
was stabilized at 5 mV/5 min. The potential was scanned
between the OCP and 300 mV/SCE in both the cathodic
and anodic directions at a scan rate of 10 mV min.–1 A
PC-controlled potentiostat (VoltaLab PGZ 301) was
applied.
3 RESULTS AND DISCUSSION
3.1 Weight-loss measurements
A weight-loss measurement for monitoring a metal’s
corrosion rate is a very usefully technique because of its
simplicity and reliability.11 The weight losses of the
copper electrodes were determined after 7 d of immer-
sion in the blank (0.1 mol dm–3 Na2SO4, pH = 3) and the
inhibitor-containing solutions. Table 1 shows the weight
loss of copper and the protection efficiency of the inve-
stigated thiazole derivative.
The protective efficiency, /%, of the DABDT mole-
cules was calculated using the following equation:12
(%) =
−
⋅
W W
W
0
0
100
where W0 and W are the weight losses of the copper
coupons in the blank and inhibitor-contains solutions.
Table 1: Results of weight-loss measurements
Tabela 1: Rezultati meritev zmanj{anja mase
c /mmol dm–3 W/mg cm–2 /%
0 1.54 –
0.001 0.80 48
0.003 0.54 65
0.005 0.25 84
0.007 0.20 87
0.01 0.14 91
It is clear from Table 1 that the corrosion rates were
reduced in the presence of the DABDT derivate. The
inhibition efficiency, as given in Table 1, is found to
increase with the increase in the concentration of the
inhibitors.
3.2 Potentiostatic measurements
Potentiostatic polarization measurements were
performed to determine the optimal concentration of
DABDT and its inhibition efficiency on copper corrosion
in acidic media (pH = 3). The polarization measurements
were in the range of the concentration 0.001 mmol dm–3
to 0.01 mmol dm–3 at the room temperature.
Figure 2 shows the anodic and cathodic polarization
plots of the copper electrode in the blank and inhibitor-
containing solution (c = 0.05 mmol dm–3).
The obtained polarization curve in the blank and the
inhibitor-containing solution are typical for copper in an
acidic solution. The cathodic part of the polarization
curves related to the oxygen reduction reaction:
O H H O22 4 4 2+ + ⇔
+ −e
and the anodic part related to the copper dissolution:
Cu Cu
Cu Cu+
⇔ +
⇔ +
+ −
+ −
e
e2
In the blank solution, at low over potentials, the
cathodic polarization curves are linear but after approxi-
\. VA[TAG et al.: THE POSSIBILITY OF COPPER CORROSION PROTECTION IN ACIDIC MEDIA ...
330 Materiali in tehnologije / Materials and technology 47 (2013) 3, 329–333
mately –300 mV/SCE the dependence begins to deviate
from linear due to diffusion phenomena, and in these
fields the measurements were not made. In the anodic
part of the curve a significant dissolution of copper is
evident. As Figure 2 shows, the presence of the DABDT
thiazole derivative induced a significant decrease in the
cathodic current density and at low over-potentials there
is an insignificantly small reduction of the anodic current
density. It can be concluded that the investigated thiazole
acts primarily as a cathodic inhibitor against copper
corrosion, hindering the oxygen reduction reaction in
acidic media. The inhibitor’s presence in the corrosion
solution does not cause a large shift of the corrosion
potential.
In Table 2 are the electrochemical parameters, such
as the corrosion potential (Ecorr) and the corrosion current
density (jcorr), which were obtained by Tafel extrapolating
the anodic and cathodic parts of the polarization curves
and the inhibitor efficiency of the DABDT molecule.
The inhibitor efficiency of the DABDT molecule was
calculated using the following equation:
(%) =
−
⋅
j j
j
0
0
100
where, j0 and j are the corrosion-current density
measured on the copper electrode in the blank and
inhibitor-containing solutions.
The obtained results are comparable with those
calculated from the weight loss measurement (Table 1).
A small difference can be observed with several
authors.13–15 This difference can be attributed to the fact
that a gravimetric measurement gives an average
corrosion rate, whereas instantaneous corrosion rates
were obtained using polarization methods.
As can be seen from Table 2 the corrosion rate in the
presence of DABDT molecules depended on the
concentration of the thiazole. The inhibitor efficiency
increased with the increase in the concentration of
thiazole. In the investigated concentration range, the
maximum concentration of 0.01 mmol dm–3 produced
the best inhibitor efficiency. A higher concentration of
inhibitor could not be tested because of the low solubility
of the DABDT. The increasing inhibitor efficiency with
the concentration indicates that DABDT molecules
protecting the copper from corrosion via the adsorption
onto the metal surface.
During the adsorption onto the metal surface the
inhibitor molecules are replacing the water molecules
that are pre-adsorbed at the metal surface. According to
Bockris,16 the adsorption of an organic molecule at the
metal/solution interface may be written according to the
following displacement reaction:
inh n inh n(sol) (ads) (ads) (sol)H O H O+ ⇔ +2 2
where n is the number of water molecules removed from
the metal surface for each molecule of adsorbed
inhibitor. The interactions of the organic molecules at
the electrical double layer change its properties and
structure. Changes in the structure of the electric double
layer occur due to the fact that the inhibitor molecules
are larger than the water molecules. Organic molecules
also have a lower value of the dielectric constant than
the water molecules, which reflects in the properties of
the double layer as the reduction of its conductivity.
Adsorption isotherms provide information about the
interaction among the adsorbed molecules themselves
and their interactions with the electrode surface.17 The
values of the degree of surface coverage, , (Table 3)
calculated from the following equation18:
 = −1
0
j
j
\. VA[TAG et al.: THE POSSIBILITY OF COPPER CORROSION PROTECTION IN ACIDIC MEDIA ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 329–333 331
Table 2: Electrochemical parameters of copper corrosion in the pre-
sence of different concentrations of DABDT
Tabela 2: Elektrokemijski parameteri korozije bakra v razli~nih kon-
centracijah DABDT
c/mmol dm–3 Ecorr/mV jcorr/μA cm–2 /%
0 –58 14.16 –
0.001 –18 7.71 46
0.003 –11 4.87 66
0.005 –20 2.74 81
0.007 –12 1.61 89
0.01 –26 1.54 90
Table 3: Copper surface coverage in a different inhibitor concen-
tration
Tabela 3: Obseg prekritja povr{ine bakra pri razli~nih koncentracijah
inhibitorja
c/mmol dm–3 
0.001 0.455
0.003 0.656
0.005 0.807
0.007 0.886
0.01 0.898
Figure 2: Polarization curves for copper electrodes with and without
the thiazole derivatives (c = 0.005 mmol dm–3)
Slika 2: Polarizacijske krivulje za bakrove elektrode brez derivatov
tiazola in z njimi (c = 0,005 mmol dm–3)
were used to determine the isotherm that best describes
the adsorption process.
The best correlation between the experimental results
and the isotherms described above was obtained using
Bockris-Swinkels isotherm:16,4
[ ]


 
( )
( )
1
1 1
−
⎡
⎣⎢
⎤
⎦⎥
=
+ − −
n
n
n
n
n
Kc
where  is the surface coverage, n is the number of
water molecules substituted, c is the inhibitor concentra-
tion and K is the constant of the adsorption process.
Figure 3 shows the Bockris-Swinkels isotherms
obtained from copper in the presence of different con-
centrations of thiazole, where f is equal to the left part of
the Bockris-Swinkels isotherms,
Knowing the value of K provides the possibility of
calculating the basic thermodynamic parameter, the
standard free energy of the adsorption is Gads. The con-
stant of adsorption, K, is related to the standard free energy
of adsorption, ΔGads, with the following equation:19
Table 4 reports the data from Bockris-Swinkels iso-
therm and the value of ΔGads for the adsorption process of
the investigated thiazole derivative.
Table 4: Value of the adsorption constant, K, the number of replace-
ment water molecules, n, the regression coefficient, R, and the stan-
dard free energy of adsorption, Gads, for the DABDT molecule
Table 4: Vrednosti adsorpcijske konstante K, {tevilo nadome{~enih
molekul vode n, regresijski koeficient R in standardna prosta energija
adsorpcije Gads za molekulo DABDT
ln K n Gads/kJ mol–1 R
13.55 1 –43.5 0.981
The high value of the adsorption constant, K, indi-
cates that the adsorption process of the DABDT mole-
cules takes place relatively quickly on the copper
surface. The results in Table 4 show that the DABDT
molecule is adsorbed on the copper surface, replacing
one molecule of water. The number of water molecules
that is replaced can be seen as an indicator of the
position of the inhibitor molecules on the surface of the
copper. The small number of water molecules suggests
that the thiazole molecules are vertically oriented in
relation to the copper surface.
The negative values of Gads ensure the spontaneity
of the adsorption process and the stability of the
adsorbed layer on the copper surface. Generally, the
values of Gads are used to determine the type of interac-
tion between the organic molecules and the metal surface
atoms. It is well known that the values of Gads of the
order of –20 kJ mol–1 or lower indicate physisorption;
those of order of –40 kJ mol–1 or higher involve charge
sharing or a coordinate type of bond between metals and
organic molecules.20,21 In case of DABDT molecules the
calculated values of Gads are slightly more negative
than –40 kJ mol–1, and therefore indicate that the ad-
sorption mechanism of this thiazole derivative on copper
in acidic sulphur-containing media was typically chemi-
sorption.
The ability to form a chemical bond with copper and
the good inhibition efficiency of the DABDT molecules
(90 %) can be explained by the characteristics of dime-
tyl-amino groups, which are present in the molecule at
position 5. In the DABDT molecule the dimethyl-amino
group has a great influence on the distribution of the
electron density around the active site of the thiazole
molecule and therefore its inhibitor activity. The dime-
tyl-amino group has a positive inductive and resonance
effect. These effects are powerful electron-donating
effects, providing an increase of the electron density of
the active centre of the thiazole molecules (sulphur
atom). According to this the unshared pair of electrons in
the sulphur atom can strongly interact with the copper
uncompleted d-orbitals to provide a protective chemi-
sorption film and better copper protection in the acidic
media.
4 CONCLUSION
In this paper the possibility of the corrosion protec-
tion of copper in acidic media using a DABDT derivative
was investigated in an acidic sulphate-containing solu-
tion (0.1 mol dm–3 Na2SO4, pH = 3) using potentiostatic
polarization measurements.
From the obtained results, it can be concluded:
1. The investigated derivative, DABDT, shows a good
inhibition efficiency (90 %) towards the copper
corrosion in the acidic sulphate-containing solution.
2. DABDT acts as a mixed copper corrosion inhibitor.
Compared to the blank solution in the presence of
these molecules there is a slowdown in the cathodic
reaction as well as the anodic process of the metal
dissolution.
3. The best inhibition efficiency was obtained for the
concentration 0.01 mmol dm–3.
\. VA[TAG et al.: THE POSSIBILITY OF COPPER CORROSION PROTECTION IN ACIDIC MEDIA ...
332 Materiali in tehnologije / Materials and technology 47 (2013) 3, 329–333
Figure 3: Bockris-Swinkels isotherms
Slika 3: Bockris-Swinkelsove izoterme
4. DABDT protects the copper surface from corrosion
by adsorption.
5. The adsorption of these molecules on the copper
surface happens as a fast and spontaneous process,
following the Bockris-Swinkels isotherms by
replacing one water molecule on the copper surface.
6. Adsorption leads to the formation of a chemical bond
between the thiazole derivative DABDT and the
copper atoms.
7. There is a good inhibition efficiency of the tested
compound achieved due to the presence of strong
electron-donating N(CH3)2 groups.
Acknowledgments
These results are the part of project No. OI-172013,
which is supported financially by the Serbian Ministry of
Education and Science.
5 REFERENCES
1 http://www.g2mtlabs.com/cost-of-corrosion/ 17. 2. 2012
2 C. Fiad, 8 SEIC, Ferrara, 10 (1995), 929–949
3 R. Fuchs-Godec,V. Dole~ek, Colloids and Surfaces A Physicoche-
mical and Engineering Aspects, 244 (2004), 73–76
4 G. Moretti, F. Guidi, Corrosion Science, 44 (2002), 1995–2011
5 R. F. V. Villamil, P. Corio, J. C. Rubim, S. M. L. Agostinho, Journal
of Electroanalytical Chemistry, 472 (1999), 112–119
6 E. S. M. Sherif, S. M. Park, Electrochim. Acta, 51 (2006),
4665–4673
7 F. Zucchi, G. Trabanelli, M. Fonsati, Corros. Sci., 38 (1996),
2019–2029
8 M. M. Antonijevic, M. B. Petrovic, Int. J. of Electrochem. Sci., 3
(2008), 1–28
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Kálmán, Solid State Ionics, 141 (2001), 87–91
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(2001), 1861–1869
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276–282
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(2006), 384–388
13 S. Muralidiharan, M. A. Quraishi, S. V. K. Lyer, Corros. Sci., 37
(1995), 1739–1750
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15 F. Bentiss, M. Lebrini, M. Lagrenee, Corros. Sci., 47 (2005),
2915–2931
16 J. O. M. Bockris, D. A. J. Swinkels, J. Electrochem.Soc., 111 (1964),
736
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19 J. Flis, T. Zakroczymski, J. Electrochem. Soc., 143 (1996),
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Acta, 53 (2008), 5941–5952
\. VA[TAG et al.: THE POSSIBILITY OF COPPER CORROSION PROTECTION IN ACIDIC MEDIA ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 329–333 333

H. JIRKOVA et al.: ENERGY- AND TIME-SAVING LOW-TEMPERATURE THERMOMECHANICAL TREATMENT ...
ENERGY- AND TIME-SAVING LOW-TEMPERATURE
THERMOMECHANICAL TREATMENT OF
LOW-CARBON PLAIN STEEL
PRIHRANKI ENERGIJE IN ^ASA PRI NIZKOTEMPERATURNI
TERMOMEHANSKI OBDELAVI MALOOGLJI^NEGA
PLO[^ATEGA JEKLA
Hana Jirkova1, Daniela Hauserova2, Ludmila Kucerova1, Bohuslav Masek1
1University of West Bohemia, The Research Centre of Forming Technology – FORTECH, Univerzitni 22, 306 14 Pilsen, Czech Republic
2COMTES FHT, a.s., Prumyslova 995, 334 41 Dobrany, Czech Republic
h.jirkova@email.cz
Prejem rokopisa – received: 2012-09-21; sprejem za objavo – accepted for publication: 2012-10-16
Reduction of energy is one way of cutting the cost of finished steel products. In this context heat treatment is a costly stage of
the production. This paper presents a method of reducing soft-annealing times. The method termed as ASR (accelerated
spheroidization and refinement) was employed for improving the cold formability of the ferrite-pearlite steel.
In this study, plain low-carbon RSt-32 steel was used as the experimental material. The influence of deformation below the Ac1
temperature on grain refinement and carbide spheroidization, as well the influence of the temperature cycling around the Ac1
temperature within the temperature intervals of various widths were explored. In addition, attention was paid to the effects of the
holding time between deformation steps and to the amount of deformation heat.
The initial lamellar pearlite was converted into a recrystallized structure with fine ferrite grains of about 3 μm and fine
spheroidized cementite. The final hardness was about 150 HV10. This is by about 25 % lower than the hardness of the initial
ferrite-lamellar pearlite structure. Significant time and energy savings can be reached as the treatment shortens the
thermomechanical exposure from several hours to several minutes.
Keywords: pearlite morphology, soft annealing, incremental deformation
Zmanj{evanje porabe energije je ena od poti za zmanj{anje stro{kov kon~nih jeklenih proizvodov. V tem kontekstu je toplotna
obdelava draga faza proizvodnje. ^lanek predstavlja metodo za zmanj{anje ~asa mehkega `arjenja. Metoda, ozna~ena z ASR
(pospe{ena sferoidizacija in zmanj{anje zrn) je bila uporabljena za izbolj{anje hladne preoblikovalnosti feritno-perlitnega jekla.
Za eksperimentalni material v tej {tudiji je bilo uporabljeno malooglji~no jeklo RSt-32. Preiskan je bil vpliv deformacije pod
temperaturo Ac1 na zmanj{anje velikosti zrn in na sferoidizacijo karbida, kot tudi vpliv nihanja temperature okrog Ac1 v razli~no
{irokih temperaturnih intervalih. Dodatna pozornost je bila posve~ena {e u~inku ~asa zadr`anja med posameznimi
deformacijami in na dele` deformacijske toplote. Za~etni lamelarni perlit se je pretvoril v rekristalizirano strukturo z drobnimi
zrni ferita, velikosti okrog 3 μm in drobno sferoidiziranim cementitom.
Kon~na trdota je bila okrog 150 HV10. To je okrog 25 % ni`je od trdote za~etne strukture s feritom in lamelarnim perlitom.
Mogo~e je dose~i pomembne prihranke v ~asu in energiji, saj termomehanska obdelava skraj{a potreben ~as od nekaj ur na
nekaj minut.
Klju~ne besede: morfologija perlita, mehko `arjenje, prirastek deformacije
1 INTRODUCTION
Reducing soft-annealing times is one of today’s
efforts in cutting down the production costs of cold-
formed parts. Initial microstructures of the materials
suitable for cold forming have a high formability. Such
microstructures can be obtained through spheroidization
of the cementite shape in pearlite.
Cementite can be spheroidized using several standard
heat-treatment procedures. These include the following:
a) isothermal annealing at a temperature slightly below
Ac1, b) soaking at a temperature just above Ac1 with slow
cooling in the furnace, or with a hold just below Ac1, c)
thermal cycling in the vicinity of Ac1 1. In this context, the
new ASR (accelerated spheroidization and refinement)
procedure is highly effective2,3 (Figure 1). It is an
energy-saving thermomechanical-treatment procedure
based on incremental deformation.
2 EXPERIMENTAL WORK
2.1 Experimental Material
The experimental material was cold-drawn, plain,
structural steel RSt37 (S232 JRC) (Table 1). The
as-received microstructure consisted of ferrite and
lamellar pearlite. The ultimate strength, yield strength,
elongation and hardness of the as-received material were
516 MPa, 450 MPa, 20 % and 200 HV10, respectively.
The pearlite fraction found with an image analysis was
9 %. The ferrite-grain size was approximately 30 μm.
The Ac1 temperature found by dilatometric testing was
777 °C.
2.2 ASR Trial
To obtain the resulting microstructure of a very fine
ferrite with spheroidized cementite, a number of process-
ing parameters had to be optimised: the strain magni-
Materiali in tehnologije / Materials and technology 47 (2013) 3, 335–339 335
UDK 669.15-194:539.377 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)335(2013)
tude, the deformation temperature and the holding time
upon deformation. The effects of deformation heat and
thermal cycling in the vicinity of Ac1 were explored as
well. The processing was performed in a thermomecha-
nical simulator.4
With the aim to reduce spheroidization times as much
as possible, pearlitic cementite lamellae must be not only
fractured but also shifted further apart (Figure 1). This is
why the schedules with two deformation steps were used
for the initial experiment (Table 2). The first step was
the tensile deformation with a strain magnitude of  =
0.3, whereas the second step included intensive com-
pressive deformation. The strain magnitude  varied
from 1.7 to 0.3 (Table 2). The effects of a 300 s hold
following the deformation upon recrystallization were
mapped. The soaking temperature and the time were
740 °C and 10 s, respectively. The heating rate was
30 °C/s. The schedules were followed by air cooling.
The effect of the thermal cycling between 740 °C and
780 °C, i.e., around Ac1, for 50 s (Table 3) was investi-
gated. Two schedules were designed: one included
deformation prior to the thermal cycling and the other
was without deformation. An additional schedule
included soaking at a reduced temperature of 700 °C and
a thermal cycling within the range of 700–780 °C for
50 s.
The effect of deformation heat on the specimen
temperature and carbide spheroidising was explored in
the subsequent schedules (Table 4). At 700 °C, the ten-
sile and compressive deformations with the magnitudes
of 0.3 and 0.4 were applied with varied holding times
between the deformation steps. The holding times were
between 5–30 s and 5–9 deformation cycles were used.
Finally, the specimens were water quenched for the
mapping of the microstructure evolution.
While the initial schedules included deformation at a
constant temperature, the additional schedules were used
to explore the effects of the non-isothermal deformation,
during which the specimen was cooled in air (Table 5).
The increase in the specimen temperature due to
deformation heat, the microstructure evolution and its
morphology, the refinement and, in particular, the
spheroidization processes were examined. A soaking
temperature of 700 °C was used in both schedules of this
stage.
H. JIRKOVA et al.: ENERGY- AND TIME-SAVING LOW-TEMPERATURE THERMOMECHANICAL TREATMENT ...
336 Materiali in tehnologije / Materials and technology 47 (2013) 3, 335–339
Figure 1: Scheme of the newly developed ASR process
Slika 1: Shematski prikaz novo razvitega postopka
Table 1: Chemical composition of the RSt37-2 (S232 JRC) steel in
mass fractions (%)
Tabela 1: Kemijska sestava jekla RSt37-2 (S232 JRC) v masnih
dele`ih (%)
C P S Mn Si N
0.08 0.022 0.023 0.65 0.16 0.004
Table 2: Thermomechanical-treatment parameters used for exploring
the effect of magnitude of compressive strain and the impact of the
hold
Tabela 2: Parametri termomehanske obdelave, ki je bila uporabljena
za iskanje velikosti tla~nih napetosti in vpliv ~asa zadr`evanja
TA/°C

Tension+
compression
Hold
upon
def. (s)
HV10
Rm/
MPa
A5mm/%
740
0.3 + 1.7 – 166 465 8
0.3 + 1.7 300 121 421 20
0.3 + 0.8 300 143 – –
0.3 + 0.3 300 141 – –
Table 3: Effect of thermal cycling on microstructure evolution
Tabela 3: Vpliv toplotnega nihanja na razvoj mikrostrukture
TA/°C

Tension +
compression
Temp.
range (°C) HV10
Rm/
MPa
A5mm/%
740 0.3 + 1.7 740–780 146 426 22
740 – 740–780 145 479 37
700 – 700–780 149 / /
Table 4: Impact of deformation heat on microstructure evolution
during a repeated deformation
Tabela 4: U~inek deformacijske toplote na razvoj mikrostrukture med
ponavljanjem deformacij
TA/°C No. of def.steps
Time between
def. steps (s) HV10
Rm/
MPa
A5mm/%
700 9 30 150 352 22
9 10 155 – –
9 5 197 260 14
5 30 140 – –
9 30 water 189 502 22
Table 5: Impact of non-isothermal deformation on microstructure
evolution
Tabela 5: Vpliv neizotermne deformacije na razvoj mikrostrukture
TA/°C No. of cycles Hold time betweendef. (s) HV10
700 9 1 225
700 9 0 192
3 RESULTS AND DISCUSSION
3.1 Effects of Strain Magnitude and Holding Time
Using the ASR schedule with the soaking tempe-
rature of 740 °C and the tensile and compressive defor-
mations ( = 0.3 + 1.7), we broke up the cementite
lamellae into the particles (Figure 1) of about 1 μm and
refined the ferrite grains. The obtained microstructure
showed a notable banding. The observation made with a
scanning electron microscope revealed that the cementite
spheroidization was extensive. The material’s hardness,
the ultimate strength and the A5mm elongation were
166 HV10, 465 MPa and 8 %, respectively (Table 2).
The elongation was rather low.
Then, an additional 300 s hold was incorporated,
following the tensile-compressive deformation, to
promote the recrystallization and to improve the
carbide-spheroidising conditions. This hold caused a
ferrite coarsening, a local carbide spheroidization and a
more uniform distribution of carbides in the ferrite
matrix (Figure 2). These microstructure changes were
reflected in a low hardness of 121 HV10 and a low
strength of 421 MPa. Recrystallization led to a higher
elongation of 20 % (Table 2).
However, the amount of the strain used is too large
for some technical applications. At the edges of the test
specimen, where the nominal strain was somewhat
lower, the ASR process appeared to be less intense.
Based on this finding, two schedules were proposed,
where the amount of strain in the second deformation
step was decreased from  = 1.7 to 0.8 and 0.3
(Table 2). The total strain levels in these schedules were
 = 1.1 and  = 0.6. The reduction in the total amount of
strain was reflected in the reduced hardness values of
143 and 141 HV10. The specimens treated with both
schedules contained large lamellar-pearlite colonies.
3.2 Effects of Thermal Cycling on Microstructure
Evolution
This stage of the study was aimed at verifying the
theory that the thermal cycling with the upper limit
above Ac1 accelerates cementite spheroidization through
a repeated dissolution of cementite and a formation of
new nuclei (Table 3). Using the schedule with the
thermal-cycling range of 740–780 °C, i.e., in the vicinity
of Ac1, the specimen microstructure consisted of the
ferrite with the grain size of above 20 μm and the pearlite
colonies along the grain boundaries with the hardness of
145 HV10 (Table 3, Figure 3). When a lower heating
temperature of 700 °C and a cycling within a wider tem-
perature range were used, the resulting microstructure
was coarser but showed the same hardness. The
incorporation of a two-step tensile-and-compressive
deformation cycle ( = 0.3 + 1.7) at 740 °C, prior to the
thermal cycling in the range of 740–80 °C led to a finer
H. JIRKOVA et al.: ENERGY- AND TIME-SAVING LOW-TEMPERATURE THERMOMECHANICAL TREATMENT ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 335–339 337
Figure 2: 740 °C/10 s – 2-times deformation step: tension + com-
pression ( = 2.1) with a hold of 300 s
Slika 2: 740 °C/10s – 2-krat stopnja deformacije: nateg + stiskanje (
= 2.1) z zadr`evanjem 300 s
Figure 4: Effects of thermal cycling in the temperature range of
740–780 °C with deformation
Slika 4: U~inek toplotnih ciklov z deformacijo v temperaturnem
podro~ju 740–780 °C
Figure 3: Effects of thermal cycling in the temperature range of
740–780 °C without deformation
Slika 3: U~inek toplotnih ciklov v temperaturnem podro~ju
740–780 °C brez deformacije
microstructure with a similar hardness value (Table 3,
Figure 4). The grains were fully recrystallized and had a
size of approximately 10 μm. In none of these cases the
cementite precipitation within the ferrite grains occurred
and the pearlite morphology was mainly lamellar.
3.3 Effects of Deformation Heat on Microstructure
Evolution
At this stage of the study, the feasibility of the
cementite spheroidization and microstructure refinement
by applying deformation heat alone was explored. This
deformation heat was to be generated by means of a
two-step deformation at 700 °C, combined with the
pauses of varying lengths (Table 4).
The schedule with 9 deformation cycles and 30 s
pauses led to an increase in the specimen temperature to
720 °C over the first 5 deformation cycles. The fine,
elongated ferrite microstructure with disintegrated pear-
lite islands and a hardness of 150 HV10 was obtained
(Figure 5). Cementite was predominantly found along
the ferrite-grain boundaries. However, there were some
particles within the ferrite grains as well. The micro-
structure was substantially refined by the forming
process. The ferrite-grain size was approximately 5 μm
and most of the cementite particles were globular. The
steel in this condition had a strength of 352 MPa and an
elongation of 22 % (Table 4). Where 10 s pauses bet-
ween deformation cycles were used instead of the 30 s
ones, the resulting extent of recrystallization was lower.
The final microstructure showed a strong banding and
disintegrated, globular pearlite islands indicating the
flow direction. The fact that the pauses were shorter by
20 s had no significant impact on the hardness. However,
reducing the pauses further to 5 s led to a large increase
in the hardness amounting to 197 HV10. The micro-
structure was similar to that obtained in the previous
case. Where the number of deformation cycles was
reduced from 9 to 5, the coarser ferrite was recrystallized
(Figure 6). Its hardness was 140 HV10 (Table 4). The
pearlite colonies disintegrated partially, remaining
lamellar in character and situated on the ferrite-grain
boundaries.
3.4 Impact of Non-Isothermal Deformation on Micro-
structure Evolution
The last series of optimisation schedules was aimed
at exploring the deformation schedules similar to the
previous ones, except that the specimen was not kept at
the heating temperature during either the deformations or
the deformation pauses (Table 5). As the specimen was
cooled in still air, the pause between deformation steps
was either 1 s or zero.
The schedule with 1 s pauses between deformations
led to a specimen failure in the 6th cycle at 520 °C. The
deformation heat was not sufficient to compensate for
the heat losses, which is why the specimen temperature
H. JIRKOVA et al.: ENERGY- AND TIME-SAVING LOW-TEMPERATURE THERMOMECHANICAL TREATMENT ...
338 Materiali in tehnologije / Materials and technology 47 (2013) 3, 335–339
Figure 6: 700 °C/10 s – 5-times tension + compression ( = 4), 0 s
pause between deformation cycles
Slika 6: 700 °C/10s – 5-krat nateg + tlak ( = 4), 0 s premora med
cikli
Figure 7: Effects of non-isothermal deformation without pauses
Slika 7: U~inek neizotermne deformacije brez premora
Figure 5: 700 °C/10 s – 9-times tension + compression ( = 7.2), 30 s
pause between deformation cycles
Slika 5: 700 °C/10 s – 9-krat nateg + tlak ( = 7.2), 30 s premor med
cikli deformacije
could not exceed 700 °C. The continuous deformation
without pauses generated a deformation heat that
increased the specimen temperature to 724 °C. The
temperature decreased below 700 °C no sooner than
during the 4th deformation cycle. In this case the speci-
men failed during the 7th cycle at 680 °C. Both schedules
resulted in highly distorted microstructures with the fine,
ferrite-grain and banded, elongated areas with globular
cementite (Figure 7). Large strains led to the hardness
values of 225 and 192 HV10.
4 CONCLUSION
The experimental programme consisted of a gradual
optimisation of the low-temperature thermomechanical
treatment of the RSt37-2 structural steel.
It was found that the processes of carbide spheroidi-
zation, grain refinement and cementite redistribution
from the bands formed upon deformation are governed
by the pause between deformation steps. Other important
preconditions for a successful ASR process include the
presence of strain components of sufficient intensity
capable of breaking up and separating fragmented
cementite particles through a plastic deformation. The
spheroidization process can be accelerated substantially
under these conditions. The optimum results of the expe-
rimental ASR process were achieved using the schedule
with the soaking temperature of 740 °C and two-step
tensile-compressive deformation with the following 300
s pause. This procedure led to a microstructure with fine,
ferrite-grain and globular carbides. The yield strength of
the resulting material was approximately 200 MPa, its
ultimate strength was 421 MPa and the elongation was
20 %.
Acknowledgements
This paper presents the results achieved in the project
GA CR P107/10/2272 "Accelerated Carbide Spheroidi-
zation and Grain Refinement in Steels". The project has
been subsidised from specific resources of the state
budget for research and development.
5 REFERENCES
1 A. Kamyabi-Gol, M. Sheikh-Amiri, Spheroidizing Kinetics and
Optimization of Heat Treatment Parameters in CK60 Steel Using
Taguchi Robust Design, Journal of Iron and Steel Research Inter-
national, 17 (2010) 4, 45–52
2 B. Masek, H. Jirkova, L. Kucerova, Rapid Spheroidization and Grain
Refinement Caused by Thermomechanical Treatment for Plain
Structural Steel, Materials Science Forum, 706–709 (2012),
2770–2775
3 D. Hauserova, J. Dlouhy, Z. Novy, J. Zrnik, M. Duchek, Forming of
C45 Steel at Critical Temperature, Procedia Engineering, 10 (2011),
2955–2960
4 B. Ma{ek et al., Physical modelling of microstructure development
during technological processes with intensive incremental defor-
mation, The Mechanical Behaviour of Materials X, 345–346 (2007),
943–946
H. JIRKOVA et al.: ENERGY- AND TIME-SAVING LOW-TEMPERATURE THERMOMECHANICAL TREATMENT ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 335–339 339

M. SHARIFIRAD et al.: ELECTROCHEMICAL CHARACTERIZATION OF THE NANO Py/DDS/SiO2 FILM ...
ELECTROCHEMICAL CHARACTERIZATION OF THE
NANO Py/DDS/SiO2 FILM ON A COPPER ELECTRODE
ELEKTROKEMIJSKA KARAKTERIZACIJA NANOPLASTI
Py/DDS/SiO2 NA BAKRENI ELEKTRODI
Meysam Sharifirad1, Farhoush Kiani2, Fardad Koohyar2
1Department of Chemistry, Teacher Research Bojnord, Iran
2Department of Chemistry, Faculty of Science, Ayatollah Amoli Branch, Islamic Azad University, Amol, Iran
fardadkoohyar@yahoo.com
Prejem rokopisa – received: 2012-09-21; sprejem za objavo – accepted for publication: 2012-11-13
The electroactive copolymer of pyrrole (Py) and 4.4’-diaminodiphenyl sulfone (DDS) was synthesized electrochemically in 4 M
H2SO4 and ethanol medium. Both electrochemical synthesis and characterization of the copolymer deposited on a copper
electrode were carried out using cyclic voltammetry. The voltammograms exhibited different patterns of behavior with different
feed concentrations of Py. Equimolar concentrations of Py and DDS demonstrated a very efficient growth of the copolymer film
on the surface of the copper. The scan rate exerted little effect on this copper copolymer film, revealing the film excellent
electroactive adherent properties. The effect of pH on the copolymer film showed that the polymer was electrochemically active
up to pH 7.0. A spectroelectrochemical analysis of the copolymer film, carried out on an indium tin oxide (ITO) plate, showed
multicolor electrochromic behavior when the applied potential was changed. The polymer was characterized with the UV-Vis
and FTIR spectral studies. The formation of the polymer through the N-H group was understood from the single N-H stretching
vibrational frequency at 3050 cm–1. The surface morphology was studied using a SEM analysis and the grain size of the
copolymer was measured using XRD studies and was found to be 50 nm. The electrical conductivity of the copolymer was
5.98 × 10–2 S cm–1, as determined using a four-probe conductivity meter.
Keywords: 4.4’-diaminodiphenyl sulfone, cyclic voltammetry, SEM, nanosize, copper electrode
Elektri~no aktiven kopolimer pirola (Py) in 4,4’-diaminodifenil sulfona (DDS) je bil sintetiziran elektrokemijsko v mediju 4 M
H2SO4 in etanola. S cikli~no voltametrijo je bila narejena elektrokemijska sinteza in karakterizacija kopolimera, nanesenega na
bakreno elektrodo. Voltamogrami prikazujejo razli~ne spektralne vzorce pri razli~nih koncentracijah dodanega Py. Ekvimolarna
koncentracija Py in DDS izkazujeta u~inkovito rast kopolimerne plasti na povr{ini bakra. Hitrost skeniranja ima majhen u~inek
na bakrovo kopolimerno plast, razkrije pa odli~ne elektroaktivne lastnosti. U~inek pH na kopolimerno plast je pokazal, da je
plast elektrokemi~no aktivna do pH 7,0. Spektroelektrokemi~na analiza kopolimerne tanke plasti na plo{~i iz indij-kositrovega
oksida (ITO) je pokazala ve~barvno elektrokromno vedenje, ~e se je uporabljeni potencial spremenil. Polimer je bil ocenjen z
UV-Vis- in FTIR- {tudijami spektra. Nastanek polimera preko N-H-skupine je razumeti kot N-H razteznostna vibracijska
frekvenca pri 3050 cm–1. Morfologija povr{ine je bila preu~evana s SEM-analizo, velikost zrn kopolimera je bila izmerjena z
XRD in ugotovljeno je bilo, da je velikosti 50 nm. Elektri~na prevodnost kopolimera, izmerjena z merilnikom prevodnosti s
{tirimi sondami, je bila 5,98 × 10–2 S cm–1.
Klju~ne besede: 4,4’-diamino difenil sulfon, cikli~na voltametrija, SEM, nanovelikost, bakrena alektroda
1 INTRODUCTION
Progressive research in the field of conducting
polymers has led to the development of the materials
with a great potential for commercial applications,
including lightweight batteries, light-emitting diodes,1
capacitors,2 electrochromic devices,3,4 optical and elec-
tronic devices.5 Conducting-polymer films have shown
promising applications in the field of biosensors and
bioelectrochemistry by providing an active matrix with
controlled morphology for immobilization of biological
materials as well as transduction of electrical signals.6,7 If
an electrode surface is modified with a conducting-
polymer film, then the modified electrode can be used as
a sensor.8–10 Polyaniline is among the most important
organic conducting polymers having a high conductivity
but poor processability.11 In order to increase its pro-
cessability and utility, researchers studied the derivatives
of polyaniline prepared with different methods. Homo-
polymerization of aniline derivatives has been effective
in the preparation of substituted polyanilines. Increased
torsional angles and the presence of a substituent result
in a decreased orbital overlap of electrons and nitrogen
lone pairs; however, substituted polyanilines exhibit
conjugations and conductivities that are significantly
lower than those of polyaniline.12 Several researchers
have employed different post treatments after the syn-
thesis of polymer films; however, there are no systematic
studies on this aspect. On the other hand, when a poly-
aniline layer was deposited first and the electropolyme-
rization was continued in a solution of 4.4’-diamino-
diphenyl sulfone (DDS), a progressive transformation of
the electrochemical behavior of the original polyaniline
film was observed. The polyaniline promotes the
polymerization by providing electrocatalytic sites and
nucleation centers. Thus, the conducting polyaniline
surface is protected well and leads to a better modifi-
cation. Chemical polymerization of DDS also leads to
newer polymers. Interestingly, both newly formed poly-
mers are found to have a nanostructure. The preparation
Materiali in tehnologije / Materials and technology 47 (2013) 3, 341–347 341
UDK 678.7:544.6 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)341(2013)
of nanostructured conducting materials has become an
important branch of materials research.13 Many research
groups have reported various kinds of polymer nanocom-
posites, like polyaniline, PEDOT, and polypyrrole.14–16
Hence, it seemed interesting to study the formation and
properties of the new nanostructured polymeric material
produced with a reaction of polyaniline and DDS during
redox cycling. DDS, an important pharmaceutical com-
pound, is used as an antileprotic drug.17–19 Utilization of
the materials of biointerest for the preparation of newer
nanopolymeric materials assumes their importance in the
present environment. They are also used in the sorption
and transport of carbon dioxide20–22 and cured in epoxy
networks.23 Here we present the results obtained from a
formation of a nanosized electroactive polymer through
an electrochemical and chemical polymerization of
4.4’-diaminodiphenyl sulfone (DDS) and the characteri-
zations of the new polymer. The electrical conductivity
of the copolymers increased greatly, from 6.00 × 10–4
S cm–1 to 2.55 × 10–1 S cm–1, with the increasing aniline
content. The UV-Vis spectroelectrochemical studies
performed on these copolymer films revealed their
electrochromic sites corresponding to individual units. In
another study an electroactive conducting polymer,
poly(pyrrole-co-4.4’-diaminodiphenyl sulfone), was
prepared from pyrrole and 4.4’-diaminodiphenyl sulfone
(DDS) using cyclic voltammetry on a copper electrode
surface.24,25 The present study reports an electrochemical
synthesis of a new nanosized poly(Py-co-DDS) and its
characterization with the SEM, XRD, and FTIR tech-
niques. Furthermore, the solubility in different solvents,
conductivity and electrochromic behavior of the
copolymer were evaluated and reported.
2 EXPERIMENTAL WORK
Pyrrole (Merck; 99 %) was distilled before use and
all the test solutions were freshly prepared. DDS and
methanol were purchased from Merck. Nano SiO2
(Allderich, Germany, 40 nm diameter particles) and a
phosphate buffer solution (PBS) with a pH of 7.0 as the
electrolyte were prepared with 0.1 M NaH2PO4–
Na2HPO4, while the pH was adjusted with 0.1 M HCl
and 0.1 M NaOH. Cyclic voltammetry (CV) was carried
out using a Potentiostat/Galvanostat EG&G Model 263
A, USA, with a PC and an electrochemical set up con-
trolled with the M 270 software. A Pt grid was utilized
as a counter-electrode and the reference electrode was
Ag/AgCl (KCl: 3 M). The working electrode was either
a copper (99.99 % purity) disk or a rectangular sheet (an
area of 0.2 cm2). The electrodes were mechanically
polished with abrasive paper (2400 grade) and rinsed
with distilled water and finally dried under an argon flow
before each electrochemical experiment. After the depo-
sition the working electrode was removed from the elec-
trolyte and rinsed with double distilled water and then
dried in air. The FTIR transmission spectrum of the
polypyrrole coating was recorded in the horizontally
attenuated, total reflectance mode in the spectral range of
400–500 cm–1 using a Bruker spectrometer, Vector Series
22, Germany. Spectroelectrochemical studies were per-
formed in a quartz cuvette with a path length of 1 cm
utilizing an optically transparent working electrode, an
indium tin oxide (ITO) plate (10  cm–2), a Pt counter
electrode, an Ag/Ag+ reference electrode and a compu-
ter-controlled JASCO V-530, UV-Vis spectrophotometer.
Scanning electron microscopy (SEM) images were taken
using a VEGA HV (high potential) 1 500 V at various
magnifications. A conducting poly(aniline-co-diamino-
diphenyl sulfone) was synthesized and characterized and
it was shown that the reactivity of 4.4’-diaminodiphenyl
sulfone was greater than that of aniline.26 Our electroche-
mically synthesized copolymer of aniline and 4.4’-dia-
minodiphenyl sulfone exhibited novel electrochromic
properties.26 In this paper, the electrochemical copolyme-
rization of pyrrole (Py) with 4.4’-diaminodiphenyl
sulfone (DDS) and nanoparticles of SiO2 is presented.
The solubility of copolymers was studied with various
organic solvents.
3 RESULTS AND DISCUSSION
3.1 Electrochemical copolymerization of pyrrole,
4.4’-diaminodiphenyl sulfone and nanoparticle
SiO2
Figure 1 shows the cyclic voltammogram of 0.01 M
DDS in 4 M H2SO4, 0.1 M SiO2 and the ethanol mixture
obtained between –1 V and 1.5 V at a scan rate of 100
mV s–1 on a copper electrode. The voltammogram exhi-
bits one broad oxidation peak at 0.2 V in the first cycle.
The anodic peak is caused by an oxidation of the amino
group in the phenyl ring of DDS in 4 M H2SO4, 0.1 M
SiO2 and ethanol medium. After the completion of the
M. SHARIFIRAD et al.: ELECTROCHEMICAL CHARACTERIZATION OF THE NANO Py/DDS/SiO2 FILM ...
342 Materiali in tehnologije / Materials and technology 47 (2013) 3, 341–347
Figure 1: Cyclic voltammetric behavior of 0.01 M DDS and SiO2 in 4
M H2SO4 and ethanol at a scan rate of 100 mV s-1 using a copper
working electrode
Slika 1: Vedenje pri cikli~ni voltametriji 0,01 M DDS in SiO2 v 4 M
H2SO4 in etanolu pri hitrosti skeniranja 100 mV s–1 z uporabo bakrene
delovne elektrode
tenth cycle, the working electrode was washed with
ultrapure water and a light brown film was seen on the
surface of the copper electrode. The film was thin and
further growth was inhibited due to its low conductivity.
Figure 2 represents the cyclic voltammograms of the
growth of PPy on a stationary copper electrode in 4 M
H2SO4, 0.1 M SiO2 and ethanol medium with a potential
range of –1 V to 1.5 V and a scan rate of 100 mV s–1. In
the first anodic scan, a peak corresponds to the oxidation
of PPy to produce a Py cation radical (PyR) that was
observed at 0.2 V. Another peak observed in this first
cycle at 1.03 V was due to the oxidation of ethanol.27
The CVs of the second and subsequent cycles during the
electrochemical polymerization of Py show two anodic
peaks at 0.3 and 0.8 V. The progressive increases in the
current of the peak at 0.3 V suggest a continuous for-
mation of PPy films on the surface of the copper elec-
trode.
The cyclic voltammogram of 0.01M DDS, 0.1 M
SiO2 and 0.01 M PPy in the 4 M sulfuric acid and
ethanol is presented in Figure 3. The potential range
applied and the scan rate selected here are the same as in
the previous cases. The CVs show two oxidation peaks at
–0.5 V and 1 V in the first cycle representing the forma-
tion of a pyrrole cation radical (PyR) and a 4.4’-diamino-
diphenylsulfone cation radical (DDSCR), respectively.
These peaks were assigned to the reduction of polymer
products formed with a reaction between the interme-
diate species PyCR and DDSCR, respectively. According
to the CVs recorded during the further cycles, i.e., the
second to the tenth cycles, one oxidation and one reduc-
tion peak at 0.8 V and 0.3 V, respectively, appeared and
they were both intensifying during these cycles. Thus,
the CVs recorded during the copolymerization of 0.01 M
DDS, 0.1 M SiO2 and 0.01 M Py in the 4 M sulfuric acid
and ethanol clearly differ from the CVs recorded during
the homopolymerization of either Py or DDS alone. The
twin redox characteristics28 noticed for the polymeri-
zation of Py were virtually merged into a single redox
process in the cases of the copolymerization of Py and
DDS. The above mentioned oxidation and reduction
waves disappeared and one anodic and one cathodic peak
appeared at 0.8 V and 0.3 V, respectively. These redox
peaks show an effective increase of the current from the
second cycle onwards. At the end of the tenth cycle, a
dark-green copolymer film was observed on the surface
of the copper electrode.
Copolymers were prepared from different molar feed
ratios of Py and the copolymer formation was influenced
by the Py concentration. The concentration of a Py mo-
nomer was plotted against the total charge of the forma-
tion of electroactive, oxidative, conducting copolymer
M. SHARIFIRAD et al.: ELECTROCHEMICAL CHARACTERIZATION OF THE NANO Py/DDS/SiO2 FILM ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 341–347 343
Figure 3: Cyclic voltammetric behavior of 0.01 M DDS, 0.01 M Py
and SiO2 in 4 M H2SO4 and ethanol at a scan rate of 100 mV s–1 using
a copper working electrode
Slika 3: Vedenje pri cikli~ni voltametriji 0,01 M DDS, 0,01 M Py in
SiO2 v 4 M H2SO4 in etanolu pri hitrosti skeniranja 100 mV s–1 z
uporabo bakrene delovne elektrode
Figure 4: Plot of the charge vs. Py concentration (obtained from the
CVs of copolymerization for varying concentrations of Py from 0.002
to 0.016 and 0.01 M of DDS in the 4 M H2SO4, 0.1 M SiO2 and
ethanol mixture)
Slika 4: Odvisnost naboja od koncentracija Py (dobljeno iz CV pri
kopolimerizaciji za razli~ne koncentracije Py od 0,002 do 0,016 in
0,01 M DDS v me{anici 4 M H2SO4, 0,1 M SiO2 in etanola)
Figure 2: Cyclic voltammetric behavior of 0.01 M Py and 0.1 M SiO2
in 4 M H2SO4 and ethanol at a scan rate of 100 mV s-1 using a copper
working electrode
Slika 2: Vedenje pri cikli~ni voltametriji 0,01 M Py in 0,1 M SiO2 v 4
M H2SO4 in etanolu pri hitrosti skeniranja 100 mV s–1 z uporabo
bakrene delovne elektrode
films (Figure 4). Figure 4 shows that the total charge is
on a rapid increase due to the increase in the Py concen-
tration until Py reaches 0.01 M, when the increase in the
charge slows down depending on the Py concentration.
These observations indicate an increasing number of Py
units in the copolymer formation.
3.2 Effect of the scan rate
The copolymer film of poly(pyrrole-co-4.4’-diamino-
diphenyl sulfone-SiO2) was washed with ultrapure water
and a monomer-free background solution and then the
film was scanned in the solution in the range of 0 mV to
1.5 mV. The cyclic voltammogram presented in Figure 5
resembles that of the electroactive polymer obtained
during the copolymerization of Py and DDS. The main
oxidation-peak current increased linearly (Figures 5 and
6) as the scan rate increased from 50 mV s–1 to 300
mV s–1 indicating the presence of the electroactive copo-
lymer film.
3.3 Effect of pH
The copolymer film formed on the surface of the
copper electrode was washed with ultrapure water and a
monomer-free solution. Then the film was scanned in the
monomer-free solution and the pH range between 4 and
9 was considered. The voltammetry range between 0 V
and 1.5 V was considered. The CV (Figure 7) resembles
that of the electroactive polymer obtained earlier. The
maximum current was observed for pH 7 and the mini-
mum for pH 4.
The peak voltammetric current decreased with the
increasing pH (Figure 7), indicating an involvement of
H+ ions. The film was also first cycled at a high pH of
7.0 (Figure 7). Thus, when the copolymer film was
scanned at pH 4 the increase in the current was due to
the change in the dopant. The effect of pH was studied
on the copolymer films of Py, SiO2 and DDS at various
monomer concentrations and similar results were
observed.
3.4 UV-Vis spectra and conductivity of copolymers
In-situ UV-Vis spectroelectrochemistry provides a
useful tool for studying the electropolymerization and
intermediate products analysis.29 To study the spectro-
electrochemical properties of the copolymer, the
copolymer film was electrochemically deposited on an
ITO glass plate at a constant potential of 1.2 V vs
Ag/Ag+. To ensure a consistent content of the electro-
active polymer on the electrode surface between poly-
M. SHARIFIRAD et al.: ELECTROCHEMICAL CHARACTERIZATION OF THE NANO Py/DDS/SiO2 FILM ...
344 Materiali in tehnologije / Materials and technology 47 (2013) 3, 341–347
Figure 7: Effect of the pH variation from 4 to 9 on the poly(Py-co-
DDS/SiO2)-film-coated copper electrode in the acidic, neutral and
basic media
Slika 7: Vpliv spreminjanja pH od 4 do 9 na poly (Py-co-DDS/SiO2)
plast, ki prekriva bakreno elektrodo, v kislem, nevtralnem in alkalnem
mediju
Figure 5: Effect of the scan-rate variation on the cyclic voltammetry
of the poly(Py-co-DDS/SiO2)-film-coated copper electrode from 50
mV s–1 to 300 mV s–1 in the 4 M H2SO4 medium
Slika 5: Vpliv spreminjanja hitrosti skeniranja od 50 mV s–1 do 300
mV s–1 v mediju 4 M H2SO4 na cikli~no voltametrijo poli (Py-co-
DDS/SiO2) plasti, ki pokriva bakreno elektrodo
Figure 6: Effect of the scan-rate variation on the poly(Py-co-DDS/
SiO2)-film-coated copper electrode from 50 mV s–1 to 300 mV s–1 in
the 4 M H2SO4 medium
Slika 6: Vpliv spreminjanja hitrosti skeniranja od 50 mV s–1 do 300
mV s–1 v 4 M H2SO4 na poly (Py-co-DDS/SiO2) plast, ki prekriva
bakreno elektrodo
merizations, the same amount of charge was passed
during copolymerization. After deposition, the blue,
oxidized ITO-adhered films were washed with the
monomer-free electrolyte solution before recording the
spectra at various applied potentials in the 0.1 M H2SO4
medium. Since all of the copolymer films are blue in
their oxidized state, each film was subsequently reduced
to determine whether there was a direct correlation
between the monomer compositions and electrochromic
response. As an illustration, the spectra of the copolymer
films obtained from 0.3 M pyrrole, 0.1 M SiO2 and 0.1
M DDS at various applied potentials are presented in
Figure 8. When the applied potentials changed from 0 V
to 1.5 V, the spectra exhibited absorption bands at 320
nm and 390 nm; the former one may be due to a -*
transition and the latter one may be a benzenoid band.
As the applied potential expanded to the oxidation side,
the film color changed from yellow to blue. Apart from
these bands, an additional broad band was observed in
the visible region. The wavelength maxima of this band
depended on the applied potentials. When an applied
potential changed from 0.0 V to 0.8 V, an absorption
band was obtained between 320 nm and 390 nm, exhi-
biting neutral yellow due to the formation of cation
radicals (polaronic forms). As the potential varied from
0.0 V to 0.8 V, the absorption band shifted to the lower-
energy side, i.e., a bath chromic shift was observed. The
non-conducting blue film may be formed because of the
fully oxidized copolymer. The feed ratio of the DDS
monomer increased to 0.5 M and the film was coated as
previously. After the deposition, the oxidation film was
dark blue, indicating a higher amount of DDS in the
copolymer. Here the -* transition and benzenoid
bands were also observed.
The conductivity of this copolymer film was measu-
red with a four-probe conductivity meter. At room tem-
perature the conductivity of the copolymer, poly(Py-co-
DDS/SiO2), was determined to be 5.98 × 10–2 S cm–1.
Thus, the conductivity of the copolymer is lower than
that of Py/SiO2 (6.68 × 10–2 S cm–1) and higher than that
of polyDDS/SiO2 (6.01 × 10–3 S cm–1).
3.5 FTIR spectral behavior of copolymers
The IR spectrum, Figure 9, reveals the presence of
different species involved in the fabrication of the
composite polymer. The bands at (3470, 3350, and 3215)
cm–1 correspond to the N-H stretching vibration, whereas
the bands at (3050, 2970, and 2850) cm–1 result from the
aromatic C-H stretching vibration. The peak at 1600
cm–1 is due to the stretching deformation of the quinone
ring. The 1280 cm–1 band is assigned to the C-N
stretching in the secondary aromatic amine, whereas the
peaks at 1080 cm–1 and 1150 cm–1 represent the aromatic
C-H in-plane bending modes. The out-of-plane deforma-
tion of C-H in the 1.4-disubstituted benzene ring is
located at 835 cm–1. The bands at 1302 cm–1 and 555
cm–1 correspond to the S=O stretching and S=O bending
mode, respectively, of the sulfone group of the PDDS
shift of their frequency and intensity in poly(Py-co-
DDS/SiO2) to 1308 cm–1 and 576 cm–1, respectively.
3.6 SEM and XRD
Chemically copolymerized materials were characte-
rized with a SEM analysis. A SEM photograph (Figure
10a) of the copolymer formed from 0.3 M Py/SiO2 and
0.02 M DDS showed a leave-like structure. This irregu-
lar structure confirmed the formation of the copolymer.
The grain size of the material is 60 nm. When the con-
centration of DDS increased to 0.03 M, the formed
copolymer exhibited only small changes in the structure
(Figure 10b). When the incorporation of DDS increased
M. SHARIFIRAD et al.: ELECTROCHEMICAL CHARACTERIZATION OF THE NANO Py/DDS/SiO2 FILM ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 341–347 345
Figure 9: FTIR spectral behavior of poly(Py-co-DDS/SiO2)
Slika 9: FTIR-spekter poly (Py-co-DDS/SiO2)
Figure 8: Spectroelectrochemical behavior of 0.3 M pyrrole, 0.1 M
DDS and 0.1 M SiO2 deposited on an ITO plate in the 0.1 M H2SO4
medium with various applied potentials
Slika 8: Spektroelektrokemi~no vedenje 0,3 M pirola, 0,1 M DDS in
0,1 M SiO2, nanesenih na ITO-plo{~i v mediju 0,1 M H2SO4 pri raz-
li~nih uporabljenih potencialih
to 0.1 M, the SEM photograph (Figure 10c) exhibited
different, irregular, broken nanostructures.
The crystalline regions in the copolymers are charac-
terized by the presence of relatively sharp peaks. The
amorphous regions are visible due to broad, low intensity
peaks. The X-ray diffraction profile (Figure 11) of the
copolymer indicates a substantial degree of crystalinity
in the doped forms. The base form of the copolymer with
low DDS/SiO2 exhibited less crystallinity than the highly
doped form. The use of the Scherrer equation is the pri-
mary technique for determining the size or thickness of
the polymer crystallites. The crystallite size of the copo-
lymer was determined by employing the XRD results
and the Scherrer equation (1):
L
K
=

  ( ) cos2
(1)
Here K is the shape factor of the average crystallite
(the expected shape factor is 0.9),  is the wavelength
(usually 0.154056 nm),  is the peak position and FW is
the full width at half maximum. By employing the above
method, the crystallite size of the copolymer was deter-
mined as 50 nm confirming the presence of nanostructu-
red copolymers. This is the first report on the crystallite
size of a copolymer of pyrrole/SiO2 and 4.4’-diamino-
diphenyl sulfone. However, earlier findings showed a
similar result for poly 4.4’-diaminodiphenyl sulfone.24
4 CONCLUSION
The copolymer of pyrrole and 4.4’-diaminodiphenyl
sulfone was prepared electrochemically using the
continuous-cycling method in a solution of 4 M H2SO4,
SiO2 and ethanol, with different concentrations of both
pyrrole and 4.4’-DDS. The copolymer formation was
more significant with the feed concentrations of 0.01 M
Py and 0.01 M 4.4’-DDS and exhibited a varying, cyclic
voltammetric behavior. The synthesized poly(Py-co-
DDS/SiO2) films demonstrated a good adherence in
acidic and neutral solvents and were found to be
electroactive. The copolymer films were responsive up to
pH 7.0 in the acid and neutral media. The formation of
the copolymer, poly(Py-co-DDS/SiO2), was further
ascertained from the FTIR and SEM data. The XRD
results revealed the formation of a nanosized copolymer.
These results indicate that the copolymers herein can be
used in the development of dynamic electrochromic
devices and their conducting films can be used to modify
the electrodes that are used as sensors.
5 REFERENCES
1 P. K. H. Ho, J. S. Kim, J. H. Burroughes, H. Becker, S. F. Y. Li, T. M.
Brown, F. Cacialli, R. H. Friend, Nature, 404 (2000), 481–484
M. SHARIFIRAD et al.: ELECTROCHEMICAL CHARACTERIZATION OF THE NANO Py/DDS/SiO2 FILM ...
346 Materiali in tehnologije / Materials and technology 47 (2013) 3, 341–347
Figure 11: XRD behavior of the poly(Py-co-DDS 0.1 M SiO2)-film
on the ITO plate
Slika 11: XRD-spekter poly(Py-co-DDS 0,1 M SiO2) plasti, ki pokri-
va ITO-plo{~o
Figure 10: SEM photographs of the copolymer (Py-co-DDS/SiO2): a) 0.3 M Py, 0.1 M SiO2 and 0.02 M DDS surface; b) 0.3 M Py, 0.1 M SiO2
and 0.03 M DDS surface; c) 0.3 M Py, 0.1 M SiO2 and 0.02 M DDS surface
Slika 10: SEM-posnetki kopolimera (Py-co-DDS/SiO2): a) povr{ina 0,3 M Py, 0,1 M SiO2 in 0,02 M DDS, b) povr{ina 0,3 M Py, 0,1 M SiO2 in
0,03 M DDS, c) povr{ina 0,3 M Py, 0,1 M SiO2 in 0,02 M DDS
2 S. Möller, C. Perlov, W. Jackson, C. Taussig, S. R. Forrest, Nature,
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M. SHARIFIRAD et al.: ELECTROCHEMICAL CHARACTERIZATION OF THE NANO Py/DDS/SiO2 FILM ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 341–347 347

P. PAV[I^ et al.: SEWAGE-SLUDGE STABILIZATION WITH BIOMASS ASH
SEWAGE-SLUDGE STABILIZATION WITH BIOMASS
ASH
STABILIZIRANJE KOMUNALNEGA MULJA S PEPELOM
BIOMASE
Primo` Pav{i~1, Danijel O{tir2, Ana Mladenovi~3, Sabina Kramar3,4,
Matej Dolenec4, Peter Bukovec5
1PMA d.o.o., Tbilisijska 61, 1000 Ljubljana, Slovenia
2VIPAP d.d., Tovarni{ka 18, 8270 Kr{ko, Slovenia
3Zavod za gradbeni{tvo Slovenije, Dimi~eva 12, 1000 Ljubljana, Slovenia
4Naravoslovnotehni{ka fakulteta, A{ker~eva 12, 1000 Ljubljana, Slovenia
5Fakulteta za kemijo in kemijsko tehnologijo, A{ker~eva 5, 1000 Ljubljana, Slovenia
primoz.pavsic@pma-lj.si
Prejem rokopisa – received: 2012-10-01; sprejem za objavo – accepted for publication: 2012-10-23
By mixing the sewage biodegradable sludge and biomass ash a stable composite is formed. A study showed that, in this way, a
building material with a compressive strength of about 2 MPa is produced, which can be used mainly for landfill covers,
road-shoulder management, road-base stabilization and rehabilitation of degraded areas. An analysis of the chemical
composition of the water eluate from the composite showed that the new composite material is inert and, as such, does not pose
a threat and does not burden the environment. From the sustainable-development point of view this kind of waste-residue
management presents an optimum – a zero waste solution.
Keywords: biodegradable sewage sludge, stabilization, biomass ash, composite construction material
Z me{anjem komunalnega biomulja in pepela biomase se tvori stabilen kompozit. Raziskave so pokazale, da tako pridobljen
gradbeni material s tla~no trdnostjo okoli 2 MPa lahko uporabimo na obmo~ju deponij komunalnih odpadkov za dnevne in
kon~ne prekrivke, za ureditev bre`in in bankin, pa tudi kot podlago za gradnjo transportnih poti in sanacijo degradiranih
obmo~ij. Tudi analiza kemi~ne sestave izlu`ka iz kompozita je pokazala, da je novi kompozit inerten in kot tak{en ne
obremenjuje okolja. Z vidika trajnostnega razvoja je tak{na re{itev optimalna, saj je rezultat predelave uporaben produkt
(predelava "zero waste").
Klju~ne besede: komunalni biomulj, stabilizacija, pepel biomase, kompozitni gradbeni material
1 INTRODUCTION
In the waste-water treatment process large amounts
of biodegradable sewage sludge are produced. Sewage
sludge, due to its high organic content, a presence of
pathogenic bacteria, heavy metals and organic pollu-
tants,1 poses a major environmental problem. Up-to-date
biodegradable sewage sludge management mainly
includes composting and reuse in agriculture, recycling
with anaerobic digestion or thermal treatment for energy
utilization and landfilling.1 Unfortunately, the reuse of
biodegradable sewage sludge in agriculture is often
impossible due to the presence of bacteria and heavy
metals, while energy utilization and landfilling still
burden the environment. In addition, the management of
different waste-ash residues is environmentally very
demanding. However, some of these products were
already successfully applied in the field of construc-
tion,2–5 though this is not true of the ashes from biomass
combustion. Alternatively, biodegradable sewage sludge
and biomass ash can be treated in the same process.6 By
mixing waste-biomass ash and biodegradable sewage
sludge, a composite construction material for specific
purposes can be obtained.7 From the sustainable-deve-
lopment point of view this presents an optimum – a zero
waste solution.
2 EXPERIMENTAL WORK
For assessing the mechanical properties of the
composite material, cube samples (150 mm × 150 mm ×
150 mm) were prepared by mixing biodegradable
sewage sludge with biomass ash cooled to ambient
temperature in the 1 : 1 mass ratio.
The compressive strength of the composite was
determined according to SIST EN 12390-3 after (7, 14,
28, 56 and 90) d of curing the specimens in moist
(100 % air moisture) and for some specimens (after 7, 28
and 56 d) also in wet conditions (submerged in water).
During the mixing process the changes in the
temperature and pH level were monitored. The eluate
was obtained from the composite after 28 days of curing
the specimens in moist conditions according to SIST EN
1744-3:2002 at the liquid/solid (L/S) ratio of 10 : 1.
Inorganic parameters of the composite eluate, biomass-
ash eluate and biodegradable sewage sludge were
determined by the inductively coupled plasma mass
spectrometry (ICP MS) according to SIST EN ISO
17294-2:2005, ISO 16772:2004-modif. and SIST EN
ISO 10304-1:2009, respectively. The effect of stabili-
zation on microbiological quality of the material was
assessed with the aerobic mesophilic bacteria count
Materiali in tehnologije / Materials and technology 47 (2013) 3, 349–352 349
UDK 69.004.8:658.567.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)349(2013)
before and after the stabilization according to the internal
method of the Pulp and Paper Institute Laboratory. For
this purpose samples were kept in aseptic conditions,
diluted in a Ringer’s solution and spread on a standard
plate count agar. After a 2 d incubation period at 37 °C, a
colony count was performed. To gather more information
on the processes accompanying the stabilization, the
mineral compositions of raw materials (sludge, ash) and
composites were also determined with the X-ray powder
diffraction (XRD) using a Philips PW 3710 diffracto-
meter and Cu K(alfa) radiation. Powdered samples were
scanned at a rate of 2°/min, over the range of 2–70°
(2Q). The results were stored on a PC and analysed by
the X’Pert HighScore Plus diffraction software. In
addition, selected samples were analysed with Fourier
transform infrared spectroscopy (FTIR), using a Perkin
Elmer Spectrum 100 spectrometer. Sixty-four signal-
averaged scans of the samples were acquired. Powder
pellets were pressed from the mixtures of the samples
with KBr at a ratio of about 1 : 200. The FTIR spectra
were recorded with a spectral resolution of 4 cm–1 in the
range of 4000–400 cm–1.
3 RESULTS AND DISCUSSION
Biodegradable sewage sludge used in this study
originates from an aerobic, biological, waste-water-treat-
ment plant at the company VIPAP d.d., where industrial
and municipal waste waters are treated. It represents a
fluid component with 0.84 % of dry matter with a pH of
7.1 and conductivity of 318 mS/m. As seen in Table 1,
concentrations of inorganic substances in the biodegra-
dable sludge, when compared to the limit values
concerning discharge of waste water,8 are not environ-
mentally problematic. The mineral composition, deter-
mined with XRD, showed that the sludge consists of
calcite, quartz, dolomite and clinochlore (Figure 1).
P. PAV[I^ et al.: SEWAGE-SLUDGE STABILIZATION WITH BIOMASS ASH
350 Materiali in tehnologije / Materials and technology 47 (2013) 3, 349–352
Table 1: Inorganic parameters of biomass-ash and composite eluates,
and of biodegradable sewage sludge
Tabela 1: Anorganski parametri izlu`kov pepela biomase in kompo-
zita, ter koncentracije preiskovanih parametrov v biomulju
Parameter Biodegradablesewage sludge Biomass ash
Composite
after 28 days
pH 7.1 12.5 11.9
Conductivity
(μS/cm) 3180 7370 18
Inorganic
parameters mg/L mg/kgs.s. mg/kgs.s.
As 0.010 <0.02 <0.02
Ba 1.2 59 2.7
Cd <0.0005 <0.005 <0.005
Total Cr 0.0035 <0.01 0.016
Cu 0.19 <0.07 0.08
Hg <0.01 <0.004 <0.01
Mo 0.0077 <0.05 <0.05
Ni 0.0050 0.048 <0.01
Pb <0.005 <0.05 0.066
2102416 Sb 0.0026 <0.006 0.018
Se <0.001 <0.01 <0.002
Zn 0.58 0.12 <0.1
Cl- 44.7 25.3 15.6
F- <1.0 4.5 <2
SO42– 50.6 <5 8.6
Note: mg/kgs.s. – milligram per kilogram of dry matter
Figure 1: X-ray diffraction pattern of biodegradable sludge (BM), biomass ash (VPPZ1) and the composites after 3, 7, 14, 28, 56 and 90 days
(VPK-1 to VPK-6, respectively)
Legend: C-calcite, T-talc, P-portlandite, Q-quartz, L-lime, G-gehlenite, CC-clinochlor, D-dolomite, Py-pyrite; CACH1: Ca8Al4O14CO2*24H2O,
CACH2: Ca4Al2O14CO9*11H2O, CaCh1H: Ca4Al2O6Cl2 *10H2O
Slika 1: Difraktogram komunalnega biomulja (BM), pepela biomase (VPPZ1) in kompozitov po 3, 7, 14, 28, 56 in 90 dneh (VPK-1 do VPK-6)
Legenda: C-kalcit, T-lojevec, P-portlandit, L-apno, G-gehlenit, CC-klinoklor, D-dolomit, Py-pirit, CACH1: Ca8Al4O14CO2*24H2O, CACH2:
Ca4Al2O14CO9*11H2O, CaCh1H: Ca4Al2O6Cl2 *10H2O
Biomass ash is a residue from a steam boiler K5 at
VIPAP d.d. According to the inorganic parameters of the
eluate (Table 1) this material is classified as a non-hazar-
dous waste.9
The ash consists of free lime (CaO), portlandite
(Ca(OH)2), gehlenite, calcite, quartz, pyrite, talc, and an
amorphous phase.
In the process of mixing the biodegradable sludge
and biomass ash a stable matrix is formed. As re-
cognized experimentally, the optimum mass ratio of
biodegradable sludge (liquid component) and biomass
ash is 1 : 1. This ratio ensures an adequate hardening
time (starting at 285 min, ending at 1140 min) and
workability of the produced material. During mixing, the
temperature is elevated up to about 45 °C and the pH
value exceeds 12. This reduced the mesophilic bacteria
count from 2.6  1011 CFU/g a. s. to 2.2 × 105 CFU/g
a. s., thus, the microbial activity was effectively inhibited
and the obtained composite product does not pose a
health threat10 when used for the intended purpose.
Furthermore, analyses of the eluate from a cube sample
of the composite after 28 days of moist curing (Table 1)
showed, that the new composite material is inert9 and, as
such, environmentally acceptable. Mechanical proper-
ties, assessed with the compressive-strength determina-
tion, are time dependent (Figure 2). The compressive
strength after 28 d ranged from 1.7 MPa (moist curing)
to 1.8 MPa (wet curing – saturated). As seen from the
compressive-strength evolution, the process of hydration
was not yet completed after 28 d. The compressive
strength after 57 d increased up to about 2 MPa or even
up to 3 MPa (saturated). Although the qualitative
mineralogical composition of the composites (Figure 1)
consisted of talc, dolomite, portlandite, calcite, gehlenite
and quartz, and did not change with time, we can
observe that the amounts of free lime and portlandite are
reduced, but the amount of calcite is increased. During
the curing period new hydration products were also
observed, namely, calcium carboaluminate (chloride)
hydrates Ca8Al4O14CO2 × 24H2O (CACH1),
Ca4Al2O14CO9 × 11H2O (CACH2) and Ca4Al2O6Cl2 ×
10H2O (CAChlH), where the CACH1/CACH2 ratio is
changing with the curing time. The formation of new
products during the curing was also confirmed with
FTIR spectra (Figure 3) that show an obvious difference
when comparing the initial materials and the composites.
This is especially indicated by the shift of bands in the
range of 1015–1075 cm–1 towards the lower wave-
numbers after the stabilisation of the composite, and by
the appearance of three additional bands at around
(3675, 3621 and 3530) cm–1 ascribed to the newly
formed hydrates.
4 CONCLUSIONS
The study showed that the stabilization process of
biodegradable sewage sludge with biomass ash in a mass
ratio of 1 : 1 effectively inhibits further microbial activity
and the associated degradation. In the process a stable
composite is formed with a 28 d compressive strength of
about 1.7 MPa. In the process of hydration the amounts
of free lime and portlandite were reduced, whereas the
amount of calcite increased due to the carbonation of the
composite. During this period, hydration products were
generated – namely, the calcium carboaluminate
hydrates. An analysis of the chemical composition of the
water eluate from the composite showed that the new
composite material is inert and, as such, does not pose a
threat for the environment. This kind of composite can
be used as a construction material mainly for landfill
covers, road-shoulder management and road-base stabi-
lization, as well as rehabilitation of degraded areas. In
view of sustainable development the described waste-
residue management presents an optimum – a zero waste
solution.
P. PAV[I^ et al.: SEWAGE-SLUDGE STABILIZATION WITH BIOMASS ASH
Materiali in tehnologije / Materials and technology 47 (2013) 3, 349–352 351
Figure 3: FTIR spectra of biodegradable sludge (BM), biomass ash
(VPPZ) and the composites after 3, 7, 14, 28, 56 and 90 days (VPK-1
to VPK-6, respectively)
Slika 3: FTIR-spekter komunalnega biomulja (BM), pepela biomase
(VPPZ) in kompozitov po 3, 7, 14, 28, 56 in 90 dneh (VPK-1 do
VPK-6)
Figure 2: Compressive strength of the composite vs. the curing time
Slika 2: Odvisnost tla~ne trdnosti kompozita od ~asa negovanja
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Soils, Volume I: Engineering Properties and Construction, Guide-
lines (IHRB Project TR-461, FHWA Project 4), Iowa Highway
Research Board, Federal Highway Administration, 2005
6 F. ^ernec, J. Zule, Stability properties of biosludge-wood ash com-
posites, International Conference Waste Management, Environmental
Geotechnology and Global Sustainable Development
(ICWMEGGSD’07 - GzO’07), Ljubljana, 2007
Available from World Wide Web: http://www.srdit.si/gzo07/papers/
97FCernec_FinalPaperGzO07.pdf
7 P. Pav{i~, Poro~ilo o rezultatih raziskav biomulja stabiliziranega z
ostanki gorenja kotla K5 v dru`bi VIPAP, PMA d.o.o., 2011
8 Uredba o emisiji snovi in toplote pri odvajanju odpadnih voda v vode
in javno kanalizacijo, Uradni list RS, 2007, 45
9 Uredba o odlaganju odpadkov na odlagali{~ih, Uradni list RS, 2011,
61
10 T. Rupel, T. Pavlica, P. Planina, J. Vidrih, S. Glas, T. Car, L. Kobola,
M. Lu{icky, D. Bo{njak, D. Sabotin, Smernice za mikrobiolo{ko
varnost `ivil, ki so namenjena kon~nemu potro{niku, IVZ RS, 2005
Available from World Wide Web: http://www.mz.gov.si/fileadmin/
mz.gov.si/pageuploads/mz_dokumenti/zakonodaja/varnost_hrane/pes
ticidi/smernice_za_mikrobiolo__ko_varnost___ivil.pdf
P. PAV[I^ et al.: SEWAGE-SLUDGE STABILIZATION WITH BIOMASS ASH
352 Materiali in tehnologije / Materials and technology 47 (2013) 3, 349–352
A. MLADENOVI] et al.: DREDGED MUD FROM THE PORT OF KOPER – CIVIL ENGINEERING APPLICATIONS
DREDGED MUD FROM THE PORT OF KOPER – CIVIL
ENGINEERING APPLICATIONS
MULJ IZ LUKE KOPER – UPORABNOST V GRADBENI[TVU
Ana Mladenovi}1, @eljko Poga~nik2, Radmila Mila~i~3, Anica Petkov{ek4,
Franka Cepak5
1Slovenian National Building and Civil Engineering Institute, Dimi~eva 12, 1000 Ljubljana, Slovenia
2Salonit Anhovo, Joint-Stock Co., Vojkova 1, 5210 Deskle, Slovenia
3Department of Environmental Sciences, Jo`ef Stefan Institute, Jamova 39, 1000 Ljubljana, Slovenia
4University of Ljubljana, Faculty of Civil Engineering and Geodesy, Jamova 2, 1000 Ljubljana, Slovenia
5Luka Koper, d. d., Vojkovo nabre`je 38, 6501 Koper, Slovenia
ana.mladenovic@zag.si
Prejem rokopisa – received: 2012-10-02; sprejem za objavo – accepted for publication: 2012-11-08
The Port of Koper, one of the biggest and the most important ports in the Northern Adriatic Sea, is constantly faced with the
problems caused by the accumulation of marine sediments inside the port, disturbing some of the port’s crucial operations.
However, these sediments can be viewed as a potential raw material and, in order to define the best way of using them in the
civil-engineering field, an extensive research project has been launched. The preliminary results of this project are presented and
discussed in the paper. So far the project has given two main results: first, the concentration of heavy metals in the aqueous
leachates is low and, secondly, in their present state, the sediments are too wet, so that there are only limited possibilities for
drying them out naturally. For this reason additional technological treatment will be needed.
Keywords: dredged mud, civil-engineering applications
Luka Koper, kot eno najpomembnej{ih pristani{~ v severnem delu Jadranskega morja, se nenehno spopada s te`avo akumulacije
sedimentov na plovnih poteh, kar povzro~a te`ave pri najbolj kriti~nih delovnih zmogljivostih pristani{~a. Po drugi strani ta
material lahko obravnavamo kot potencialno surovino v gradbeni{tvu. V prispevku so podani preliminarni rezultati
interdisciplinarnih raziskav, ki ka`ejo naslednje: prvi~, koncentracija te`kih kovin v izlu`kih je nizka in drugi~, v stanju, kot je,
je sediment preve~ vla`en, da bi ga bilo mogo~e osu{evati z naravnimi postopki in je zato potrebna dodatna tehnolo{ka
obdelava.
Klju~ne besede: mulj, uporaba v nizkih gradnjah
1 INTRODUCTION
The Port of Koper is one of the biggest and the most
important ports in the Northern Adriatic Sea, and it is
primarily transit oriented. It is a multi-purpose port with
two piers, 26 berths, and 12 specialized terminals. One
of its main problems is related to a constant accumu-
lation of marine sediments inside various parts of the
port, resulting in disturbances to some of its most crucial
operational capacities. A total of 80000 m3 of sediments
have to be removed annually. The sediment found in the
Port of Koper is a mixture of clay and silt (henceforth
referred to as mud) and represents the kind of waste, for
which there is insufficient disposal space along the
Slovenian coast (Figure 1). According to the slogan "No
waste here, just resources!" this sediment can be viewed
as a potential raw material. In order to define the best
way of using it in civil-engineering applications, an
extensive research project is under way. Some of the
preliminary results are presented and discussed in the
paper.
2 EXPERIMENTAL WORK
2.1 Materials and methods
In January 2012 several batches of mud were taken
from the temporary deposits at the end of Pier I (the mud
from Basin 1) and Pier II (the mud from Basin 2). All the
material was homogenised and divided into subsamples
for the targeted analyses – except for the chemical analy-
ses, which were performed on the mud from each loca-
tion separately and, additionally, on the sample from
Basin 3. Quantitative mineralogical compositions were
determined using X-ray diffraction (Phillips PANalytical
Materiali in tehnologije / Materials and technology 47 (2013) 3, 353–356 353
UDK 69.004.8:658.567.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)353(2013)
Figure 1: The mud in the temporary landfill at Koper Port
Slika 1: Mulj na za~asni deponiji v Luki Koper
X’Pert PRO equipment with Cu K
 radiation). Powdered
samples were scanned at a rate of 2 °/min, over the range
of 2–70 ° (2). A particle-size analysis was performed
using a Cilas 920 laser granulometer in the water as a
suspension media. The samples for the thermal testing
were heated in cylindrical pots, made of inert materials,
by applying a temperature scan of 10 °C/min at room
temperature, up to 1400 °C. A STD 2960 Simultaneous
DTA-TGA analyser was used, complemented with the
Universal Analysis for Windows 95/98/NT software tool,
edition 2.5H.
Several geomechanical tests were performed in order
to evaluate the basic properties and behaviour of the mud
under loaded conditions. The most crucial test was the
edometer test carried out according to SIST/ISO/TS
17892-5:2004.
In order to evaluate the environmental impact, the
extent of pollution was estimated by determining the
total metal concentrations, by identifying the most hazar-
dous, highly mobile metal fractions and by partitioning
the metals into easily, moderate soluble, and sparingly
soluble fractions. For this purpose various sequential
extraction procedures were applied.1–3 Special attention
was paid to the estimation of the content of the highly
mobile, biologically available fractions of metals in the
sediments. The total metal concentrations in the sedi-
ments and in various stages of the sequential-extraction
procedure were determined using inductively coupled
plasma mass spectrometry (ICP-MS) on an Agilent
7700x ICP-MS. A CEM Corporation CEM MARS 5
Microwave Acceleration Reaction System was used to
digest the sediments. Mechanical shaking of the samples
was performed with a Vibromix 40 elliptical, orbital
shaker. The samples were centrifuged using a Hettich
Universal 320 Centrifuge, and a WTW 330 pH meter
was used to determine the pH values.
2.2 Reagents
Merck suprapur acids and Milli-Q water (Direct-Q 5
Merck suprapur acids and ultrapure water) were used for
the preparation of the samples and standard solutions.
All the other reagents were of the analytical reagent
grade. A Stock IV CertiPUR ICP Multi Element Stan-
dard Solution containing (1000 ± 10) mg L–1 element
concentrations in 1 mol L–1 HNO3 was obtained from
Merck. Sartorius 0.45 μm cellulose nitrate membrane
filters with 25 mm diameters were used in the filtration
procedure.
2.3 Determination of the total metal and aqueous
leachate concentrations
0.2 g of a dry homogenised sediment was weighed in
a Teflon tube, and microwave-assisted digestion using a
mixture of HNO3, HCl and HF was applied according to
the procedure proposed by [~an~ar et al.3 All the analy-
ses were performed in triplicates. In order to determine
the element concentrations of the aqueous leachates,
10 g of each sample were shaken for 24 h with 100 mL
of water, centrifuged and filtered through 45 μm
membrane filters. The elements in the aqueous leachates
were determined by ICP-MS.
3 RESULTS AND DISCUSION
3.1 Particle-size analysis
The results of the particle-size analysis showed that
the majority of the grains had a dimension of less than
2 mm, and that the quantity of the grains with a size of
63 μm or less varied between 75 % and 97 % by mass.
About 40 % by mass of the grains were under 5 μm, i.e.,
within the clay size range.
3.2 Mineralogical composition
The quantitative mineral composition of an average
sample of the mud determined by XRD is presented in
Table 1. The results show that clay minerals, i.e., illite,
chlorite and Ca montmorillonite, made up more than
40 % of the investigated mud. These results are in good
correlation with those obtained by Ogorelec4.
Table 1: Quantitative mineral composition of the mean sample of mud
Tabela 1: Kvantitativna mineralna sestava povpre~nega vzorca mulja
Mineral % by mass
Illite/muscovite 25
Chlorite 20
Quartz 21
Calcite 19
Feldspar 9
Dolomite 3
Pyrite 2
Ca montmorillonite 1
3.3 Thermal analysis
Within the range between 200 °C and 450 °C a strong
exothermic peak was observed. This peak represents an
oxidation of organic matter and dewatering of clayey
mineral assemblages and probably some amorphous
Fe-Al oxide gels (Figure 2). Within the temperature
interval between 400 °C and 500 °C the S + O2  SO2
reaction takes place, which is due to the decomposition
of pyrite into pyrrhotite and free oxygen, which reacts at
the same time with H2+; H2 + 1/2O2  H2O within the
temperature range of 530–580 °C, producing another
exothermic peak.5 At a temperature of about 580 °C the

   inversion of quartz and a likely dehydroxylation
reaction of kaolinite6 occur. In the same endothermic
interval, illite loses its constitutive water, which is caused
by dehydroxilation of the octahedral sheet, as the hydro-
xyl groups of the tetrahedral sheet are gradually removed
up to a temperature of 850 °C.7 The second endothermic
peak at about 600 °C and another, smaller one, at around
700 °C can be attributed to chlorites.3 The third and
A. MLADENOVI] et al.: DREDGED MUD FROM THE PORT OF KOPER – CIVIL ENGINEERING APPLICATIONS
354 Materiali in tehnologije / Materials and technology 47 (2013) 3, 353–356
fourth endothermic peaks form a decarbonatization inter-
val: (a) calcium carbonate derived from marine fauna
skeletons, i.e., a biogenic carbonate; (b) dolomite; and
(c) terrigenic carbonate.8 The final exothermic peak, at
approximately 800 °C, belongs to various phase trans-
formations of amorphous Fe-Al minerals.6
3.4 Geomechanical parameters
The initial research was focused on determining the
basic physical and mechanical parameters that are
crucial for evaluating a possible stabilization of the mud.
It was concluded that the content of the water in the mud
is extremely variable (from 55 % to 95 % by mass) and
that the density of the untreated material is very low so
that it is not possible to stabilise it in its existing form. It
can be seen, from the edometer curves, that, at an effec-
tive normal pressure of 200 kPa, it is possible to achieve
a pore coefficient of between 0.9 and 1.1, which corres-
ponds to the moisture content of between 33 % and 40.7
%. Only at such moisture contents (i.e., between 33 %
and 40 %) the material is close to the state when a
chemical stabilization by means of inorganic binders,
e.g., fly ash, lime, cement and their various combina-
tions, is technically and technologically feasible.
3.5 Environmental analysis
In order to estimate the extent of pollution, the total
element concentrations and the concentrations in the
aqueous leachates were determined. From the results of
these analyses and a comparison with the limiting values
prescribed by the existing legislation for the inert and
non-hazardous waste, it is clear that the highest total
element concentrations were found in the sediment from
Basin 3. The elevated total element concentrations were
observed in the case of Cr (within the range from 200
mg/kg to 380 mg/kg), Ni (within the range from 140
mg/kg to 175 mg/kg) and As (within the range from 15
mg/kg to 40 mg/kg), i.e., the concentrations that are
higher than those found in the sediments elsewhere in the
Slovenian costal area.3 In general, the concentrations in
the aqueous leachates are also the highest for the
sediment from Basin 3, but still lower than the limits set
by the legislation for the inert waste leachates.9 In order
to estimate the partitioning of the elements between the
easily and sparingly soluble sediment fractions, a 6-step
sequential-extraction procedure was also performed. The
data obtained indicated that the measured concentration
of the elements in the easily soluble fractions was, in
general, very low. The only exception was Mo. The
partitioning of Mo is presented in Figure 3. It is clear
that the easily soluble fraction (i.e., the water-soluble
fraction and the exchangeable fraction) makes up about
25 % of the sediment from Basin 1 and about 10 % of
the sediment from Basins 2 and 3. It can be concluded
that, with regard to the total heavy-metal concentrations,
the concentrations in the aqueous leachates and the
partitioning of the elements between easily and sparingly
soluble sediment fractions, the sediments from the basins
of the Koper Port can be used as a secondary material in
civil engineering.
4 CONCLUSIONS
Preliminary results show that the mud from the Port
of Koper does not represent a threat to the environment
with regard to heavy-metal pollution. However, due to its
salt content it is not possible to plan the use of this mud
in the areas outside the Port of Koper. It is planned to use
an appropriate treatment to fix the chlorides within the
building composite. Due to a high quantity of water in
the mud, it will be more economical to stabilize it with a
drying-out process, using suitable technology.
Acknowledgments
The results were obtained within the ARRS project
"Sediments in aquatic environments: their geochemical
and mineralogical characterization, remediation and use
as secondary raw materials", 2011–2014.
A. MLADENOVI] et al.: DREDGED MUD FROM THE PORT OF KOPER – CIVIL ENGINEERING APPLICATIONS
Materiali in tehnologije / Materials and technology 47 (2013) 3, 353–356 355
Figure 2: TGA-DTA analysis
Slika 2: TGA-DTA analiza
Figure 3: Partitioning of Mo between various phases (I: Water-
soluble, II: Exchangeable, III: Bound to carbonates, IV: Bound to Fe
and Mn oxides, V: Bound to organic matter, VI: Residual fraction) in
sediments (B1, B2 and B3) from Luka Koper
Slika 3: Porazdelitev Mo med razli~ne faze sedimenta (B1, B2 in B3)
iz luke Koper: (I: vodotopni, II: izmenljivi, III: vezan na karbonate,
IV: vezan na Fe in Mn okside, V: vezan na organsko snov, VI: te`ko
topni ostanek)
5 REFERENCES
1 A. [ömen Joksi~, S. A. Katz, M. Horvat, R. Mila~i~, Comparison of
single and sequential extraction procedures for assessing metal
leaching from dredged costal sediments, Water, Air and Soil
Pollution, 162 (2005), 265–283
2 R. Mila~i~, J. [~an~ar, S. Murko, D. Kocman, M. Horvat, A complex
investigation of the extent of pollution in sediments of the Sava
River, Part 1, Selected elements, Environmental Monitoring and
Assessment, 163 (2010), 263–275
3 J. [~an~ar, T. Zuliani, T. Turk, R. Mila~i~, Organotin compounds and
selected metals in the marine environment of Northern Adriatic Sea,
Environmental Monitoring and Assessment, 127 (2007), 271–282
4 B. Ogorelec, M. Mi{i~, J. Faganeli, P. Stegnar, B. Vri{er, A. Vukovi~,
The recent sediment of the Bay of Koper (Northern Adriatic), Geolo-
gija, 30 (1987), 87–121
5 G. Locher, Mathematical models for the cement clinker burning pro-
cess, ZKG International, 55 (2002), 29–50
6 R. C. Mackenzie, The differential thermal investigation of clays,
Mineralogical Society (Clay Minerals Group), London 1957, 456
7 J. H. Araújo, N. F. da Silva, W. Acchar, U. U. Gomes, Thermal
decomposition of illite, Material Research, 7 (2004), 359–361
8 @. Poga~nik, J. Pav{i~, A. Meden, Geological records as an indicator
of the thermal characteristics of mudstones within the temperature
range of decarbonatisation, Mater. Tehnol., 43 (2009) 3, 157–163
9 Official Gazette of the Republic of Slovenia, No. 61/11, 2011
A. MLADENOVI] et al.: DREDGED MUD FROM THE PORT OF KOPER – CIVIL ENGINEERING APPLICATIONS
356 Materiali in tehnologije / Materials and technology 47 (2013) 3, 353–356
P. MARUSCHAK et al.: AUTOMATED DIAGNOSTICS OF DAMAGE TO AN ALUMINUM ALLOY ...
AUTOMATED DIAGNOSTICS OF DAMAGE TO AN
ALUMINUM ALLOY UNDER THE CONDITIONS OF
HIGH-CYCLE FATIGUE
AVTOMATIZIRANA DIAGNOSTIKA PO[KODBE ALUMINIJEVE
ZLITINE PRI VISOKO-CIKLI^NEM UTRUJANJU
Pavlo Maruschak1, Igor Konovalenko1, Mykhailo Karuskevich2, Vladimir Gliha3,
Toma` Vuherer3
1Ternopil Ivan Pul’uj National Technical University, 46001 Ternopil, Ukraine
2National Aviation University, Aircraft Design Department, Komarova Ave. 1, 03058 Kyiv, Ukraine
3University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia
maruschak.tu.edu@gmail.com
Prejem rokopisa – received: 2012-10-02; sprejem za objavo – accepted for publication: 2012-11-12
An identification and quantitative analysis of the deformation relief of the aluminium alloy for an aircraft construction based on
a digital-image processing has been performed. The behaviour of defects has been assessed on the basis of diagnostics results
for individual stages of the deformation process. It has been established that the individual stages of the damage-accumulation
process are characterised by the values of integral-image parameters. Based on the consecutive processing of the data on the
surface cyclic deformation, the main regularities of the propagation of defects have been found. Theoretical preconditions have
been substantiated and experimental results obtained.
Keywords: fatigue, surface, digital image, diagnostics, accumulated damage, defect propagation, evaluation
Identificirali in kvantitativno analizirali smo relief deformirane povr{ine aluminijeve zlitine za gradnjo letal. Temelj obojega je
izdelava digitalne podobe povr{ine vzorcev. Oceno vedenja defektov omogo~ajo rezultati diagnostike posameznih stopenj
deformacijskega procesa. Ugotovili smo, da so te stopnje med akumulacijo po{kodbe zna~ilne po vrednostih integralnih
parametrov slike. Na osnovi zaporedne obdelave podatkov cikli~ne deformacije na povr{ini smo ugotovili glavne zakonitosti pri
{irjenju defektov. Tako smo dokazali teoreti~ne predpostavke in pridobili eksperimentalne podatke.
Klju~ne besede: utrujanje, povr{ina, digitalna podoba, diagnostika, akumulirana po{kodba, {irjenje defekta, vrednotenje
1 INTRODUCTION
An analysis of the loading conditions of modern civil
aircrafts, the existing methods for evaluating the accu-
mulated fatigue damage, the peculiarities of fatigue
damage of aviation structural materials and the results of
the previous fatigue investigations allowed formulating
an approach to solving the problem of a quantitative
evaluation of the accumulated fatigue damage of the
aircraft structural elements1. However, the technological
complexity of many of the existing instrumental methods
for evaluating the accumulated fatigue damage as well as
their insufficient accuracy and reliability limit the use of
these methods for practical purposes.2
The initial diagnostics of the contemporary aircraft-
skin condition involves the search and identification of
the fatigue damage using the visual-control methods. It
is known that the incubation period of the fatigue
damage accumulation is, in many cases, reflected in the
visual signs, which determine the possibility of both
qualitative and quantitative evaluations of the accumu-
lated fatigue damage.3 The quantitative evaluation of the
accumulated damage at the initial stage of fatigue allows
predicting the place and the time of a fatigue-crack
appearance. At the stage of designing aviation equipment
such a prediction reduces the cost of a full-scale fatigue
tests significantly due to shortening their duration, and,
at the stage of operation, it allows increasing the reliabi-
lity of aircrafts and safety of flights.4
A deformation relief is formed on the surface of the
cladding layer of aluminium alloys under the stresses
corresponding to the loading conditions of many struc-
tural elements during operation and testing. Stress con-
centration causes a prior formation of the relief in the
vicinity of riveting holes, glue-cooking points, etc.,
which are the areas of potential failures. A surface defor-
mation relief is observed at several scale levels. Using
the optical microscopy the signs of a relief can be
observed at the meso- and macro-levels.5
The need for substantiating and implementing the
objective indicators of the deformation-relief intensity as
the characteristics of the accumulated fatigue damage is
obvious. The solution of this problem, by means of an
analysis of damaged-surface images, will be shown
below. The dependencies that allow predicting the
residual life are of the most practical importance. Such
dependences can be obtained on the basis of the data on
deformation-relief-parameter evolution.6
It is proposed herein to use the integral parameters
obtained by analysing the investigated surface images.
Materiali in tehnologije / Materials and technology 47 (2013) 3, 357–361 357
UDK 620.178.3:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)357(2013)
They allow evaluating the fatigue damage of the alumi-
nium alloy for the aircraft construction.
2 DEFORMATION-RELIEF-EVALUATION
TECHNIQUE
The specimen geometry is shown in Figure 1; the
specimen was tested with cantilever bending at R = 0,
max = 147 MPa. The analysis of the deformation relief
was performed near the stress concentrator (a hole with a
diameter of 1.0 mm).
In the process of testing we recorded the surface
condition in the vicinity of the stress concentrator after
an application of cyclic loading, the number of loading
cycles necessary for the initiation of a fatigue crack with
a length of 1.0 mm, and the number of loading cycles
leading to a complete failure. The specimen surface con-
dition was evaluated by analysing its photo images
obtained in a series after a certain number of loading
cycles. The images taken with a photo camera (Figures
1b to d) were transformed into grayscale images with the
brightness function I(x,y). The absolute values of hori-
zontal and vertical gradients were calculated for every
pixel of an image:
∇ =I x y
I x y
xx
( , )
( , )∂
∂
, ∇ =I x y
I x y
yy
( , )
( , )∂
∂
(1)
where x N m∈ , y N n∈ (m and n are the width and
height of an image, respectively):
The mean values of horizontal and vertical gradients
were used for a generalised evaluation of the condition
of the surface investigated:
G I
mn
I x y x y
G I
mn
I x y x y
x x x
mn
y y y
= ∇ = ∇
= ∇ = ∇
∫∫
1
1
11
( , )
( , )
d d
d d
11
mn
∫∫
(2)
The gradient allows determining the predominant
direction of the defect propagation and the nonunifor-
mity degree of the surface investigated.7 A low mean
value of the gradient indicates an insignificant variation
of intensities along the given axis of an image. In
practice this shows a more uniform picture of the
deformation relief in a certain direction8 and indicates
the coordinate axis that corresponds to the predominant
direction of the defect propagation. In order to enhance
the informative features corresponding to the elements of
the damaged surface, binary transformation was applied
to the obtained grayscale images.9 This resulted in
black-and-white images of the damaged surface with the
intensity function IB, in which white pixels correspond
to the background and black ones to the objects of the
deformation relief (Figures 1e to g). The most general
parameter that allows evaluating the degree of specimen
damage using the obtained images, is the relative area of
defects:
S
S
m nd
=
⋅
⋅100% (3)
where S is the number of deformation relief pixels in an
IB image.
The distribution of deformation relief elements along
the image axes is described with horizontal Hx and
vertical Hy histograms:8
H y I x yx
x
m
( ) ( , )=
=
∑
1
, H x I x yy
y
n
( ) ( , )=
=
∑
1
(4)
Each element of a histogram contains a number of
pixels that correspond to the objects of the deformation
relief, in columns and lines of the image analysed,
respectively. Histogram functions (4) contain the basic
information about the distribution of the deformation
relief along the coordinate axes of an image.
For a generalised evaluation of the surface damage
based on histograms it is proposed in8 to use the mean
values of histograms  x S n= / and  y S m= / (where S
is the general number of black pixels). However, it is
reasonable to use these parameters during multiple
measurements under similar conditions with a permanent
rectangular watch window. During laboratory testing of
different specimens, especially under different con-
ditions of the surface defect nucleation, the mean values
of histograms contain little information. In addition, in
the case of a rectangular watch window the values of μx
and μy are scaled differently (relative to the image
P. MARUSCHAK et al.: AUTOMATED DIAGNOSTICS OF DAMAGE TO AN ALUMINUM ALLOY ...
358 Materiali in tehnologije / Materials and technology 47 (2013) 3, 357–361
Figure 1: a) Scheme of the specimen investigated; examples of the
deformation relief after 15, 100, 711 thousand loading cycles: b), c),
d) initial image and e), f), g) binary image
Slika 1: a) Shema vzorca za preiskavo; primer reliefa deformirane
povr{ine po 15, 100, 711 tiso~ obremenilnih ciklih: b), c), d) za~etna
slika in e), f), g) binarna slika
dimensions) and are inconvenient for a comparison,
while in the case of a square watch window they become
similar.
For a quantitative evaluation of the histogram view
(4) a spectral analysis of the functions was performed.
Using the fast Fourier transformation the histogram
functions were presented in the form of a row:
H y A
k
n
y
H x A
k
n
x
x xk x
k
K
y yk y
x
( ) cos( )
( ) cos( )
≈ −
≈ −
=
∑ 2
2
0
π
π


k
K y
=
∑
0
(5)
The number of the harmonics of histograms Kx and
Ky was chosen in such a way as to ensure that the accu-
racy of presenting a histogram function as a sum of
harmonics is not lower than the limit value of :
H y A
k
n
y
H x A
k
n
x
x xk x
k
K
y yk
x
( ) cos( )
( ) cos(
≈ − ≤
≈ −
=
∑ 2
2
0
π
π
 
y
k
K y
)
=
∑ ≤
0

(6)
The mean amplitudes of the spectrum of the func-
tions of the horizontal Aax and vertical Aay histograms
were taken as the informative parameters:
A
K
Aax
x
xk
k
K x
=
− =
∑1
1 1
, A
K
Aay
y
yk
k
K y
=
− =
∑1
1 1
(7)
The mean amplitude of the spectrum yields the
quantitative evaluation of the damage propagation along
the image axes. Its higher values correspond to a higher
degree of damage along the given axis. Thus, while
comparing the values of Aax and Aay it is possible to get
information about the size and predominant direction of
the surface-defect propagation.
The presence of the pairs of generalized characte-
ristics – the mean gradients Gx, Gy and the mean ampli-
tudes of the spectrum Aax, Aay – allows obtaining the
complex integrated characteristics of the image analysed
in two, mutually perpendicular, coordinate directions.
3 REGULARITIES IN DAMAGE
ACCUMULATION
Cyclic loading forms a deformation relief on the
surface of structural aluminium alloys, the intensity of
which indicates the level of the accumulated fatigue
damage.10 The relief of this type was observed on both
the standard specimens for the fatigue tests in a broad
range of loading conditions and the specimens prepared
from the skin of the An-24 aircraft; both were tested
under stresses close to the operational ones.1–3 The
results of the investigations carried out using the
methods of optical and electronic microscopy show the
appropriateness of using the deformation-relief term and
applying it as a diagnostic parameter of fatigue
damage.1,2
Figure 2 shows the dependence of the calculated
integral image parameters on the number of loading
cycles. An increase in the Sd parameter indicates an
increase in the degree of damage to the surface
investigated. According to the experimental data, in the
P. MARUSCHAK et al.: AUTOMATED DIAGNOSTICS OF DAMAGE TO AN ALUMINUM ALLOY ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 357–361 359
Figure 2: Dependence of integral parameters of deformation relief on:
a) cyclic loading of a specimen, N: general area of defects Sd, b) mean
horizontal and vertical gradients Gx and Gy, c) mean amplitudes of the
spectrum, Aax and Aay
Slika 2: Odvisnost integralnih parametrov reliefa utrujenosti od: a)
cikli~ne obremenitve vzorca N: skupna povr{ina defektov Sd, b) pov-
pre~en horizontalni in vertikalni gradient Gx and Gy, c) srednja
amplituda spektra Aax and Aay
cases of up to 100,000 loading cycles, fast growing
surface defects take place (Figure 2a). At the same time,
after 100,000 loading cycles the deformation relief
changes insignificantly.
The quantitative analysis of the kinetics of the
changes in the deformation-relief orientation during
cyclic loading of the aluminium alloy is performed. It is
found that cyclic loading of 80,000 cycles causes a for-
mation of a complex system of shears on the specimen
surface. At this stage, an intense saturation of the surface
with deformation shears takes place. Individual crystal
blocks and their conglomerates sometimes ascend and
sometimes descend above the surface, while their forma-
tions coalesce in groups and the sizes of defects increase
gradually. This dynamic chaos is a depiction of the
dynamic, cooperative dislocation processes at the micro-
and meso-levels.1,5 Nearly identical values of the gra-
dients during the first phase of the investigation reflect
the chaotic disorderly process of defect nucleation.
Later on the relief acquires an orderly orientation.
Lower values of the vertical gradient (as compared to the
horizontal one) show that along the vertical axis the
picture of the deformation relief is more uniform, while
along the horizontal axis sharper changes in the image
intensity are observed. Thus, lower values of the gradient
correspond to the direction, along which the areas of the
surface damage stretch. A monotonous increase in the Gx
and Gy parameters confirms the results obtained earlier,11
indicating self-similarity and scaling of the formed de-
formation structures. The mean amplitudes of the spec-
trum of histogram functions Aax and Aay (Figure 2c)
duplicate, to a certain extent, the dependence presented
in Figure 2a; however, they allow characterising the
surface failure process in two coordinate directions.
Higher values of the vertical amplitude indicate a more
significant damage in this particular direction.
The quantitative parameter of fatigue failure is found
and it allows assessing the surface condition and the
defect propagation direction within the section analysed
based on the analysis of the surface saturation with
visual signs of the deformation relief (Figure 3).
The chosen diagnostic parameter of damage and the
technique of its quantitative evaluation allow considering
the developed method as the express diagnostic of
fatigue failure.
4 GRADED NATURE OF THE
DEFORMATION-RELIEF DEVELOPMENT
The first stage is the fast accumulation of defects on
the surface analysed. The structural non-uniformity
causes a formation of residual stress fields, which relax
partly by means of forming sliding bands, extrusions and
intrusions in the individual grains of the material12.
Characteristic structural traces of the localisation are
observed in macro extrusions in the form of the aggre-
gates of "ridges" ("hills"). An intense formation of the
deformation relief takes place. This stage accounts for
approximately 15–20 % of the general number of cycles
to the onset of crack formation.
The second stage is the slowing down of microplastic
shears of the grains of the material, coalescence of indi-
vidual sections of the deformation relief and saturation
of the system. Although the damage accumulation pro-
cess (deformation relief) is slowed down at this stage as
compared to the previous one, its course is quasistatio-
nary. It is this that allows using it for technical diagno-
stics of a damaged-surface condition with a view to
predicting the limit state. The limit state refers to the
nucleation of a fatigue crack with the length of 1.0 mm
that can be identified by means of optical control. The
limiting value of damage parameters was considered as
the last one in the row of values that were obtained
before the moment of the fatigue-crack formation.
5 CONCLUSIONS
A technique has been developed for evaluating the
aluminium-alloy surface condition by analysing its
image and calculating its integral parameters, including
the general area of damage, the mean gradients along the
coordinate axes, and the mean amplitude of the spectrum
of histogram functions along the coordinate axes. The
basic factor that allows using the proposed technique is
P. MARUSCHAK et al.: AUTOMATED DIAGNOSTICS OF DAMAGE TO AN ALUMINUM ALLOY ...
360 Materiali in tehnologije / Materials and technology 47 (2013) 3, 357–361
Figure 3: a) Horizontal and b) vertical histograms for depicting the
deformed surface after 15, 100, 711 thousand loading cycles
Slika 3: a) Horizontalni in b) vertikalni histogram za opisovanje
deformirane povr{ine po 15, 100, 711 tiso~ obremenilnih ciklih
the saturation of the material surface with the visual
signs of the relief that are detected with the methods of
optical microscopy. For the quantitative characteristic of
the relief, the damage parameters can be used and they
allow analysing the variation of spatial orientation of the
defects caused by a deformation within the surface, on
which the signs of localised strain are absent.
The main regularities in the accumulation of fatigue
damage on the surface of an aluminium alloy have been
established. The possibility of using the developed tech-
nique for determining the accumulated fatigue damage
under high-cycle fatigue is substantiated.
6 REFERENCES
1 M. Karuskevich, O. Karuskevich, T. Maslak, S. Schepak, Int. J. of
Fatigue, 39 (2012), 116
2 S. R. Ignatovich, Materials Science, 47 (2011), 636
3 Y. G. Gordienko, R. G. Gontareva, J. S. Schreiber, E. E. Zasimchuk,
I. K. Zasimchuk, Adv. Eng. Mater., 8 (2006), 957
4 M. V. Karuskevich, E. Yu. Korchuk, A. S. Yakushenko, T. P. Maslak,
Strength of Mat., 40 (2008), 693
5 L. S. Derevyagina, V. E. Panin, A. I. Gordienko, Physical Mesome-
chanics, 11 (2008), 51
6 E. E. Zasimchuk, M. V. Karuskevich, A. I. Radchenko, Strength of
Materials, 22 (1990), 1855
7 P. V. Yasnii, I. V. Konovalenko, P. O. Marushchak, Materials Science,
45 (2009), 291
8 A. Hassani, H. Ghasemzadeh Tehrani, Crack detection and classi-
fication in asphalt pavement using image processing, in Pavement
Cracking: Mechanisms, Modelling, Detection, Testing and Case
Histories, CRC Press, Chicago 2008, 891–896
9 I. V. Konovalenko, P. O. Marushchak, Optoelectronics, Instrumen-
tation and Data Processing, 47 (2011), 360
10 P. Yasniy, P. Maruschak, I. Konovalenko, Strain, 47, (2011), 238
11 M. V. Karuskevych, I. M. Zhuravel’, T. P. Maslak, Materials Science,
47 (2011), 621
12 Yu. G. Kabaldin, S. N. Murav’yev, Russian Engineering Research,
27 (2007), 513
P. MARUSCHAK et al.: AUTOMATED DIAGNOSTICS OF DAMAGE TO AN ALUMINUM ALLOY ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 357–361 361

K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
WELDED ALUMINIUM AND MAGNESIUM ALLOYS –
CORROSION AND MECHANICAL PROPERTIES FOR
REFRIGERATION COMPRESSORS IN COMPARISON
WITH DEEP-DRAWING STEEL
VARJENE ALUMINIJEVE IN MAGNEZIJEVE ZLITINE –
KOROZIJSKE IN MEHANSKE LASTNOSTI ZA KOMPRESORJE
HLADILNIKOV V PRIMERJAVI Z JEKLOM ZA GLOBOKO
VLE^ENJE
Klaus Günther Kerschbaumer, Rudolf Vallant, Norbert Enzinger, Christof Sommitsch
Institute for Materials Science and Welding, Kopernikusgasse 24, 8010 Graz, Austria
rudolf.vallant@tugraz.at
Prejem rokopisa – received: 2012-10-03; sprejem za objavo – accepted for publication: 2012-11-12
Increasing the energy efficiency of household refrigeration appliances as a result of legal requirements is more and more
important (2010/30/EUEG2010 Directive). This means that the equipment with the energy efficiency lower than D can no
longer find its place in the European trading. In this investigation aluminium alloys (AW5083-O, AW6181-T4) and a
magnesium alloy (AZ31) were selected via a material selection. They are compared to the currently used deep-drawing steels
(DD11, DD13) with respect to the corrosion and strength properties of similar overlap-welded joints.
To verify the corrosion properties the neutral salt spray test (NSS) and the fruit acid spray test were performed with an overall
test duration of 480 h. The type of corrosion, its influence on the corrosion rate and the strength of the welded joints were
evaluated. Magnesium shows a very high corrosion and therefore cannot be used uncoated, like deep-drawing steel. The
aluminium alloys show only slight selective corrosion phenomena and are, from the welding and corrosion point of view, an
attractive alternative to steel. Due to a higher thermal conductivity of aluminium, in comparison with steel, a higher energy
efficiency of the cooling compressor is expected.
Keywords: MIG/CMT-P welding, aluminium, magnesium, corrosion, deep-drawing steel, fruit acid, neutral salt spray test,
tensile test, AW5083-O, AW6181-T4, AZ31, DD11, DD13, DIN 8985, DIN ISO 6227
Vedno bolj postaja pomembno pove~anje energijske u~inkovitosti gospodinjskih hladilnih naprav kot posledica legalnih zahtev
(2010/30/EUEG2010 Directive). To pomeni, da oprema z energijsko u~inkovitostjo, manj{o od D nima ve~ mesta v evropski
trgovini. V tej raziskavi sta bili pri izbiri materiala izbrani aluminijevi zlitini (AW5083-O, AW6181-T4) in magnezijeva zlitina
(AZ31). Primerjane so s sedaj uporabljanima jekloma za globoko vle~enje (DD11, DD13) glede na korozijo in trdnostne
lastnosti podobnih zvarov s prekrivanjem.
Za preverjanje korozijskih lastnosti sta bila izvr{ena preizkusa v nevtralni slani komori (NSS) in preizkus s {kropljenjem sadne
kisline v povpre~nem trajanju 480 h. Na zvarjenih spojih je bila ocenjena vrsta korozije, vpliv na hitrost korozije in na trdnost
zvarjenih spojev. Magnezij izkazuje veliko korozivnost in ga zato ni mogo~e uporabljati brez povr{inske za{~ite tako kot jeklo
za globoki vlek. Aluminijevi zlitini izkazujeta samo rahle selektivne korozijske pojave in sta s stali{~a varjenja in korozije
zanimivi kot alternativa za jeklo. Zaradi ve~je toplotne prevodnosti aluminija v primerjavi z jeklom se pri~akuje tudi ve~ja
energijska u~inkovitost kompresorja za hlajenje.
Klju~ne besede: MIG/CMT-P-varjenje, aluminij, magnezij, korozija, jeklo za globoki vlek, sadna kislina, nevtralni preizkus
{kropljenja s slanico, natezni preizkus, AW5083-O, AW6181-T4, AZ31, DD11, DD13, DIN 8985, DIN ISO 6227
1 INTRODUCTION
The main focus of the technological development of
household cooling compressors is on an "energy-effi-
ciency improvement" stimulated by the increased compe-
tition, cost pressures and a stricter EU legislation1,2 for
cooling compressors. Furthermore, the noise reduction in
service and an improvement of the corrosion properties
are important aims for a long-term system improvement
of household cooling compressors. These future require-
ments can be met, in part, by finding a new, innovative
housing material.
Currently, the steel materials DD11 and DD13 are
used for deep-drawn, almost spherical housings.3 The
main problem is, however, that the material corrosion
resistance in the above mentioned application is not
satisfactory. Preliminary studies have shown that alumi-
num and magnesium alloys have a potential to replace
the steel materials. For a material-selection process,
according to Reuter4, the following housing requirements
have been defined:
• the minimum static strength of 85 MPa for the maxi-
mum internal pressure of 40 bar
• MIG or equivalent weldability
• the materials should be industrially formable (cold or
hot forming)
• hermetic-gas proofness
The Al-alloys AW5083-O (annealed) and AW6181-
T4 (naturally aged) and the magnesium alloy AZ31 were
selected as suitable candidates for further investigations.
Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 363
UDK 669.715:669.721.5:620.193 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)363(2013)
Corrosion investigations of these base materials and their
similar welded joints in the neutral salt spray (NSS) and
fruit acid spray tests (FS) have been performed.
2 EXPERIMENTAL WORK
2.1 Base materials and weld filler metals
MIG pulsed-arc welding was used for the annealed,
AlMgMn, wrought alloy AW5083-0 and the naturally
aged, AlMgSi, wrought alloy AW6181-T4. For the
deep-drawing steels DD11 and DD13, MAG pulsed-arc
welding was applied. The CMT+P welding procedure
was applied for magnesium AZ31. The sheet thickness
of the materials was 3 mm, except for the 2.5 mm
AW6181-T4.
The filler metals were selected according to the
recommendations of DVS 09135 and Davis.6 The AlMg5
filler was selected for AW5083-O, and AlSi5 for
AW6181-T4. Considering the Böhler Welding Guide for
DD11 and DD13, the G3Si1 (EMK6) filler was used7.
For the AZ31 sheet, Kammer8 recommends AZ61.
AM509 was used as it easily meets the strength require-
ments of 85 MPa.
It is notable that the AlSi5 filler metal has only about
50 % of the strength of the AW6181-T4 sheet. Never-
theless, it was used as the hot-cracking susceptibility of
6000 alloys can be significantly reduced according to
Davis.6 The 440 MPa yield strength of the G3Si1 weld
metal is much higher than the 170 MPa yield strength of
deep-drawing steels, as there is no even matching filler
available. The chemical compositions and mechanical
properties of the examined sheet alloys and filler metals
are summarized in Table 1.3,7,9–12
2.2 Experimental set-up and investigations
2.2.1 Welding experiments
In Figure 1 the joint geometry and welding torch
position are shown schematically for different material
combinations. The top and bottom sheets were
positioned with an overlap of 20 mm and clamped on the
welding table. The welding experiments were carried out
with a Fronius TPS4000+CMT welding machine. The
torch was conducted with an ABB robot IRB140. The
torch angle  varied between 45° and 50° depending on
the material combinations. The contact tip distance
K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
364 Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377
Table 1: Chemical analyses and mechanical properties of sheet alloys (GW) and weld-filler metals (FM)
Tabela 1: Kemijska analiza in mehanske lastnosti plo~evin (GW) ter varilnega dodajnega materiala (FM)
material
BM / FM
element
AW5083-O
(BM)10
AlMg5
(FM) 11
6181-T4
(BM) 10
AlSi5
(FM) 11
DD11
(BM) 3
DD13
(BM) 3
EMK 6
(FM) 7
AZ31
(BM) 12
AM50A
(FM) 9
Si 0.4 <0.25 0.7–1.1 4.5–6.0 0.9 0.1
Fe 0.4 <0.4 0.15–0.5 <0.6 balance balance balance 0.008
Cu 0.1 <0.1 0.25 <0.3 0.01
Mn 0.4–1.0 0.05–0.2 0.4 <0.15 0.6 0.40 1.49 0.2 0.26–0.60
Mg 4.0–4.9 4.5–5.5 0.6–1.0 <0.2 balance balance
Cr 0.05–0.25 0.05–0.2 0.15 -
Zn 0.25 <0.1 0.3 <0.1 1.0 0.22
Ti 0.15 0.06–0.2 0.15 <0.15
Be - <0.0003 0.1 <0.0003
others 0.15 <0.05 0.15 <0.05
Al balance 3 4.4-5.4
C 0.12 0.08 0.08
P 0.045 0.03
S 0.045 0.03
Rp0.2/MPa 145 110 125 40 170 170 440 170 125
Rm/MPa 300 240 235 120 <440 <400 530 240 230
A5/% 22 17 23 8 28 33 30 17 10
Figure 1: Welding-torch positions for similar overlap welded joints
Slika 1: Pozicija varilnika pri enako debelih prekrivnih varjenih spojih
covers from 12 mm to 15 mm. The contact point of the
welding wire at the lower sheet for all the material
combinations is about 1.5 mm away from the fillet weld
root.
For all the welding experiments optimised parameters
were used to get good wetting at the weld flank. For
AW5083-O and AW6181-T4 the welding parameters
refer to Kerschbaumer13. The fillet welds of deep-draw-
ing steels were welded utilising the optimised parameter
sets of the company ACC Austria for the 30 mm/s weld
speed. For the Al und Mg experiments the welding speed
of 15 mm/s was selected. The forehand angle 
 varied
between 0° (AW5083-O, AZ31) and 5° (AW6181-T4,
DD11 and DD13). The shielding gases of 50 : 50 Ar and
He and Corgon 18 (which is 18 : 82 CO2 and Ar) were
used for Al, Mg and steel welding. The energy input
varied between 0.9 kJ/cm (AZ31) and 3.4 kJ/cm (DD13).
Aluminium and DD11 were welded with a similar
energy input. All the welding parameters are listed in
Table 2.
2.2.2 Tensile and corrosion tests
From the welded overlap joints the 50 mm × 180 mm
samples for corrosion and shear tests as well as metallo-
graphic investigations were prepared and the 50 mm ×
100 mm ones were prepared from the base materials,
Figure 2. The shear strength of the joints was tested
according to EN ISO 9018:200614 utilizing a Zwick
RMC100 tensile-testing machine. The free clamping
length was adjusted to 120 mm. The strain rate (traverse
speed) was 2 mm/min at room temperature. For com-
parability, the measured fracture forces were converted
to the tensile strength (50 mm sample width, 3 mm or
2.5 mm sheet thickness). The arithmetic mean and the
standard deviation have been calculated for each sample.
The corrosion resistance of the welded joints and
base materials was investigated using the neutral salt
spray test (NSS) according to EN ISO 9227:200615 as
well as the fruit acid spray test (FS) according to DIN
8985.16 Corrosion testing was performed using a Köhler
corrosion chamber HKT500. The samples were taken out
after (2, 6, 24, 48, 96, 168, 240 and 480) h, documented
macroscopically with a digital reflex camera Nikon D50
and the mass gain was determined with a scale of the
company Denver Instrument MXX-612 with an accuracy
of 0.01 g, according to DIN 50905-1.17 Due to a limited
number of the samples the determination of the mass
loss after removing the corrosion product was not carried
out. After 240 h and 480 h metallographic investigations
with LOM (Zeiss Z1m) and REM (Zeiss 1415VP) and
the tensile-shear tests were performed to detect the
impact of corrosion on the weld strength.
K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 365
Table 2: Welding parameters for the selected material combinations/TPS 4000/CMT
Tabela 2: Parametri varjenja za izbrane kombinacije materialov /TPS 4000/CMT
Base material AW5083-O AW6181-T4 DD11 DD13 AZ31
Filler metal AlMg5 AlSi5 EMK6 EMK6 AM50
Current type Pulse Pulse Pulse Standard CMT+P
Sheet thickness (mm) 3 2.5 3 3 3
Shielding gas/quantity
(L/min)
50 : 50 Ar:He
17
50 : 50 Ar:He
17
Corgon 18
18
Corgon 18
18
50 : 50 Ar:He
17
rapdefaultWire feed drive
(m/min) 7.1 5.8 7.5 11.5 11
Voltage U/V 19 24.6 23.3 29.4 12.3
Current I/A 123 113 244 344 113
Welding speed v/(mm/s) 15 15 30 30 15
Energy input
E/(kJ/cm)
1.55 1.85 1.90 3.37 0.93
Comments 
 = 0°
 = 50°
LBK -2

 = 5° forehand
 = 50°

 = 5° forehand
 = 50°

 = 5° forehand
 = 50°
LBK-1

 = 0°
 = 50°
Figure 2: Corrosion samples; a) welding joint for tensile-shear test
and metallography; b) base material
Slika 2: Vzorci za korozijo; a) varjeni spoj za natezni stri`ni preizkus
in metalografijo; b) osnovni material
3 RESULTS
3.1 Theoretical background
3.1.1 Deep-drawing steel
In saline solutions or brines, uniform- or shallow-
pitting corrosion occurs, and occasionally there is also
pitting corrosion, if chloride is deposited locally.18–20 In
the seawater with a salt concentration of about 29.2 g/l
shallow pitting or pitting corrosion appears on deep-
drawing steel. In the hydrochloric acid the corrosion rate
increases linearly with the temperature and acid concen-
tration.18,20 In organic acids iron is corrosion resistant as
long as the access to oxygen is prohibited. Generally, the
uniform corrosion attack increases with an increase in
the organic-acid concentration.18,20
3.1.2 Aluminium
In the atmosphere the less noble aluminium (with an
electrochemical potential of –1.66 V) forms an Al2O3
dense oxide layer. This oxide layer causes a chemical
resistance in the pH range of 4.5 to 8.8. For lower and
higher pH values, i.e., for acidic or alkaline attacks the
passive layer is dissolved and a uniform corrosion takes
place. In chloride media aluminium shows pitting,
intergranular corrosion and stress-corrosion cracking. In
drinking water aluminium and Al2O3 are resistant unless
the Al2O3 layer is mechanically damaged, e.g., due to
grinding.18,19,21–25 In milk aluminium forms an oxide layer
with a good resistance. It is attacked by salt, hydrofluoric
acid and alkali.18
Aballe has found in his work that precipitates
[Al(Mn, Fe, Cr) and Al(Si, Mg)] cause the pitting in
AW5083-O. The precipitates Al(Mn, Fe, Cr) are more
noble than the base material matrix and the corrosion
takes place in the Al matrix.26,27
An intergranular corrosion attack is formed along the
Mg2Si precipitates at the grain boundaries in
AW61814-T4. The following precipitates stimulate an
intergranular corrosion attack: CuAl2 in 2xxx28, Mg2Al3
in 5xxx25, MgZn2 in 7xxx25 and Mg2Si29 or CuAl228 in
6xxx. There are two types of intergranular corrosion,
dependent on the nobility of the grain boundaries. The
less noble Al8Mg5 precipitates in the 5xxx alloys30,31 and
the MgZn2 precipitates in the 7xxx alloys are dissolved
anodically. With the more noble grain-boundary
precipitates like Mg2Si or CuAl2, the surrounding matrix
material is dissolved.18,19,22,24,32–35
3.1.3 Magnesium
In the chloride-containing atmospheres and in salt
solutions magnesium is not resident, as the protective
oxide layer MgO is dissolved. In natural brines magne-
sium-salt crystals are formed that dissolve easily in
water. Therefore, no further protective oxide layer is
formed because in an aqueous chloride solution the
passive layer is dissolved and no other corrosion
protection is in place. In neutral media like humid air,
MgO is transformed into a protective Mg(OH)2 layer2.
Manganese and zinc have a positive effect on the
corrosion resistance of magnesium, where the Zn content
must be below 3 %. When alloying aluminium and
magnesium, the corrosion resistance in salt water
increases18,20,21,30,33,36–38.
3.2 Corrosion behaviour in the salt spray and fruit
acid tests
Macroscopically a distinct difference in corrosion
behaviour between the welded and base-material sheets
is visible if the NSS and FS tests are compared. The steel
sheets show a relatively thick rust layer after a 480 h
NSS test, whereas after a FS test only a slight rust layer
can be seen. It is noticeable that during the FS test rust
layers are formed only on the welded steel joints, not on
their base sheets. In the case of magnesium pitting can
be seen macroscopically after the NSS and FS tests. In
the case of aluminium both AW5083-O and AW6181-T4
show formations of white corrosion products (aluminum
hydroxide) during the NSS and FS tests and they are
more distinctive for the latter (Figures 3a and b).
3.2.1 Salt spray test (NSS) – metallographic
investigations
3.2.1.1 Base materials
After 480 h, on the AW5083-O surface, a few small
pitting sites could be detected with a stereo microscope
(Figure 4a), while AW6181-T4 showed no pitting. Both
alloys were covered with white corrosion products that
were more pronounced on AW6181-T4, Figure 4b. The
steel and magnesium sheets showed strong corrosion
K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
366 Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377
Figure 3: Corrosion samples after 480 h: NSS test (left), FS test
(right)
Slika 3: Vzorci s korozijo po 480 h: NSS-preizkus (levo), FS-preizkus
(desno)
K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 367
Figure 5: Corrosion after the NSS test/cross-section details in LOM: a) AW5083-O, 240 h; b) AW5083-O, 480 h; c) AW6181-T4, 240 h; d)
AW6181-T4, 480 h
Slika 5: Korozija po NSS-preizkusu; detajli prereza, svetlobna mikroskopija (LOM): a) AW5083-O, 240 h; b) AW5083-O, 480 h; c)AW6181-T4,
240 h; d) Aw6181-T4 po 480 h
Figure 4: Corrosion after the NSS test/surfaces shown with a stereo microscope: a) AW5083-O, 480 h, b) AW6181-T4, 480 h; c) AZ31, 240 h; d)
DD11, 240 h
Slika 4: Korozija po NSS-preizkusu; povr{ine, prikazane s stereomikroskopom: a) AW5083-O, 480 h, b) AW6181-T4, 480 h; c) AZ31, 240 h; d)
DD11, 240 h
already after 240 h. On the AZ31 sheet, a porous white
corrosion product was formed, Figure 4c. The steels
show abundant red rust, whereat the first signs of
corrosion occurred already after a few hours, Figure 4d.
The cross sections confirm the pitting corrosion for
the AW5083-O alloy observed after 168 h with the stereo
microscope. Significant pits were found after 240 h.
Figure 5a shows a pit depth of about 0.16 mm observed
with a light optical microscope (LOM). Primary precipi-
tates and impurities containing Fe, Si and Mg were
detected at the bottom of the pit with an energy-disper-
sive X-ray analysis (EDX) using a scanning electron
microscope (SEM), Figures 6a and b. These precipitates
behave electrochemically nobler and get exposed to the
corrosion. The pitting progress into the depth from 240 h
to 480 h is marginal (Figure 5b). The AW6181-T4 alloy
shows intergranular corrosion (IG) starting in the small
pitting areas and reaching a depth of about 0.1 mm
(Figures 5c and d). The corrosion progress from 240 h
to 480 h shows no further increase in the depth but
spreading. The first superficial corrosion attack was
visible already after 2 h. On closer examination, primary
K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
368 Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377
Figure 6: SEM and EDX of Figure 5: a) and b) AW5083-O, detail 6.1; c) and d) AW6181-T4, detail 6.2
Slika 6: SEM- in EDX-posnetka podro~ij na sliki 5: a) in b) AW5083-O, detajl 6.1; c) in d) AW6181-T4, detajl 6.2
Figure 7: Corrosion after 480-h NSS/base-material cross sections in LOM: a) DD11 – uniform corrosion and shallow pitting; b) AZ31 grinded
sample – uniform corrosion; c) AZ31 non-grinded sample – severe corrosion attack
Slika 7: Korozija po 480 h NSS; prerez osnovnega materiala, svetlobna mikroskopija (LOM): a) DD11 – enakomerna korozija in plitve jamice;
b) AZ31 bru{eno – enakomerna korozija; c) AZ31 nebru{eno – resen pojav korozije
precipitates could be detected. In the SEM investigations
no secondary phases (T4 aged) could be found, Figures
6c and d.
The cross section of the DD11 steel shows uniform
corrosion and shallow pitting. It can be expected that
DD13 has a similar corrosion behaviour, Figure 7a.
The magnesium alloy AZ31 exhibits a strong depen-
dence on the initial surface roughness. Generally, the
grinded samples have uniform corrosion and shallow
pits, which are then fortified in depth. But the non-grin-
ded AZ31 samples show a very strong local-corrosion
attack combined with uniform corrosion. After 480 h
almost the entire 3-mm sheet was dissolved, i.e.,
corroded as seen in Figures 7b and c. This significant
difference can be attributed to the effect that a corrosive
medium adheres less on the surfaces with low roughness.
3.2.1.2 Welded joints
In general, no effect of the used welding parameters
on the corrosion behaviour was detected, compare13. The
corrosion attack after 480 h was analyzed in the cross
sections of the fillet welds, Figure 8.
Pitting corrosion occurred at the heat affected zone
(HAZ) of alloy AW5083-O and in the AlMg5 weld
metal. At the weld edges and the fusion line, local
corrosion was detected, too. Furthermore, the surface
pores in the AlMg5 weld metal were found to be the
starting points for pitting and crevice corrosion as seen
on Figure 9, compare.13
Sanchez-Amaya39 also found this corrosion beha-
viour with the corrosion tests involving a mixture of
NaCl and H2O2. In the investigations of Mutombo in
201140, AW5083-O was MIG-welded with both filler
metals, AlMg5 and AlSi5, to a Y-joint and tested in 3 %
NaCl. He traced the corrosion start back to the existing
surface pores. The tough AlMg5 weld and HAZ showed
pitting.
The AW6181-T4 welded joints showed no corrosion
at the fusion line or on the AlSi5 weld metal.
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Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 369
Figure 8: AW5083-O welded joint/fillet-weld cross section in LOM
prior to the NSS test; hydrofluoric acid/alcoholic etchant
Slika 8: AW5083-O varjeni spoj; prerez zvara, svetlobna mikroskopija
LOM pred NSS-preizkusom; kislina HF in alkoholno jedkalo
Figure 9: Welded AW5083-O joint cross sections in LOM/corrosion after the 480 h NSS test/local corrosion at the fusion line and weld-metal
surface pores: a)–d) details 9.1 and 9.2 from Figure 8
Slika 9: Svetlobna mikroskopija (LOM) pre~nega prereza zvara AW5083-O; korozija po 480 h NSS-preizkusa; lokalna korozija na liniji taljenja
in pore na povr{ini materiala zvara: a)–d) detajli 9.1 in 9.2 s slike 8
The welded steel joints showed an intense corrosion
attack along the fusion line and selective corrosion of the
G3Si1weld metal, Figures 10a to d.
The AZ31 welded joints were manufactured only
from the grinded magnesium alloy and were NSS tested.
Figures 11a and b shows strong and partially localized
corrosion on the AM50 weld metal as well as on the base
material. Obviously, the bottoms of the formed cavities
(large pits) act as starters for more pitting corrosion
progressing into the depth.
3.2.2 Fruit acid test (FS) – metallographic
investigations
3.2.2.1 Base materials
For AW5083-O, pitting corrosion of about 0.1 mm in
depth occurred after the 240 h FS test. Continuing the
corrosion test to 480 h, the pits did not progress in depth
but grew wider, to about 0.6 mm, Figures 12a and b.
This was caused by an electrolytic dissolution of metal
in an acidic medium, compare.10 In contrast, AW6181-T4
shows intergranular corrosion after 240 h (0.15 mm
deep), just slightly progressing in depth (0.18 mm)
during the 480 h FS test. However, the IG corrosion
spread by about 2–3 mm in width after 480 h, Figures
12c and d.
The DD11-steel base material exhibits pitting
corrosion after the 240 h and 480 h FS tests. In the
etched cross section, partial IG corrosion could be found,
too, Figures 13a to c. The results of the corrosion
K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
370 Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377
Figure 11: Welded AZ31 joint, grinded – cross sections in LOM/
corrosion after the 480 h NSS test: a) lower weld edge; b) upper weld
edge – strong corrosion of the AM50 weld metals and the base metal
Slika 11: Zvarjen AZ31 spoj, bru{eno – pre~ni prerez, svetlobna
mikroskopija LOM; korozija po 480 h NSS-preizkusa: a) spodnji rob
zvara, b) zgornji rob zvara – mo~na korozija AM50 zvarjene kovine in
osnovnega materiala
Figure 10: Welded steel-joint details – cross sections in LOM/corrosion after the 480 h NSS test/local corrosion at the fusion line: a) DD11; b)
DD13; c and d) G3Si1 weld metal
Slika 10: Detajli zvarjenega spoja jekla – pre~ni prerez s svetlobno mikroskopijo (LOM); korozija po 480 h NSS-preizkusa; lokalna korozija na
liniji taljenja: A) DD11; b) DD13; c) in d) G3Si zvarjena kovina
behaviour of the DD13-steel base material are quite
similar to the DD11 results.
The grinded AZ31 magnesium base material shows
shallow pitting after the 240 h FS. Contrary to the NSS,
during the 480 h test the corrosion progresses strongly,
whereat localized corrosion and pitting arise, especially
at the backside of the sample, Figures 14a and b.
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Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 371
Figure 14: Corrosion of AZ31 after the FS test/cross-section details in
LOM: a) 240 h; b) 480 h
Slika 14: Korozija AZ31 po FS-preizkusu; detajli pre~nega prereza,
svetlobna mikroskopija (LOM): a) 240 h; b) 480 h
Figure 13: Corrosion of DD11 steel after the FS test/cross-section
details in LOM: a) 240 h; b) 480 h; c) 480 h, etched with Nital
Slika 13: Korozija DD11 jekla po FS-preizkusu; detajli pre~nega
prereza, svetlobna mikroskopija (LOM): a) 240 h; b) 480 h; c) 480 h,
jedkano z nitalom
Figure 12: Corrosion after the FS test/cross-section details in LOM: a) AW5083-O, 240 h; b) AW5083-O, 480 h; c) AW6181-T4, 240 h; d)
AW6181-T4, 480 h
Slika 12: Korozija po FS-preizkusu; detajli pre~nega prereza s svetlobno mikroskopijo (LOM): a) AW5083-O, 240 h; b) AW5083-O, 480 h; c)
AW6181-T4, 240 h; d) AW6181-T4, 480 h
3.2.2.2 Welded joints
After the 240 h FS test, alloy AW5083-O exhibits
pitting corrosion at the fusion line and on the base
material. After 480 h distinct pitting is found in the
AlMg5 weld metal, too. At the weld edges a strong
localized corrosion attack is found along the marked
fusion line, whereat the corrosion depth remained to be
about 0.8 mm between 240 h and 480 h, Figures 15a
to c.
The welding joint AW6181-T4 does not lead to the
corrosion of the weld edges or the AlSi5 weld metal.
Here the IG corrosion is found on the base material as
shown in Figures 12c and d.
After the 240 h FS test, the welded DD11-steel joint
shows local corrosion at the weld edges. The
shallow-pitting corrosion in the G3Si1 weld metal can be
seen after 480 h, Figures 16a and b. These results are
similar to the ones for the NSS test, Figure 10a.
In the AM50 weld metal of the AZ31 magnesium
joint, a notable amount of pores appeared. They had no
significant influence on the corrosion behaviour during
the FS test. Figures 17a and b shows shallow pitting
corrosion as well as crevice corrosion at the cold weld
edge after 240 h. These results are less critical than those
related to the strong localized corrosion in the NSS test.
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372 Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377
Figure 15: Welded AW5083-O joint cross sections in LOM/corrosion
after the FS test: a) 240 h; b) 480 h, lower weld edge; c) 480 h, upper
weld edge (see Figure 8)
Slika 15: Pre~ni prerez zvarjenega spoja AW5083-O, svetlobna
mikroskopija (LOM); korozija po FS-preizkusu: a) 240 h; b) 480 h,
spodnji rob zvara; c) 480 h, zgornji rob zvara (glej sliko 8)
Figure 17: Welded AZ31 joint, grinded – AM50 weld metal details/
corrosion after the 240 h NSS test: a) lower weld seam; b) upper weld
seam
Slika 17: Spoj zvarjenega AZ31, bru{eno – detajl AM50 zvara;
korozija po 240 h NSS-preizkusa: a) spodnji {iv zvara; b) zgornji {iv
zvara
Figure 16: Welded DD11-steel joint details – corrosion after the FS
test: a) 240 h; b) 480 h, G3Si1 weld metal
Slika 16: Detajli zvarjenega spoja DD11-jekla; korozija po
FS-preizkusu; a) 240 h; b) 480 h G3Si1 zvar
3.3 Corrosion influence on the welded-joint strength
The tensile-shear test results before and after the
corrosion testing are shown in Figures 18a and b. The
fracture forces were recorded and standardized with the
initial cross-sectional area (50 mm × 3 mm and 50 mm ×
2.5 mm for AW6181-T4). The fracture location for all
the welded shear- test samples runs along the fusion line.
In the as-welded condition, the DD11 and DD13 steels
had the highest weld strength (352 MPa and 298 MPa).
The difference is due to the lower carbon content of
DD13. AW5083-O samples show a weld strength of 274
MPa and AW6181-T4 a strength of 113 MPa. Magne-
sium AZ31 has the lowest weld strength that is 104 MPa.
Considering the influence of the corrosion time and
corrosive medium, the following can be observed:
NSS test
The DD13-steel weld showed a 12 % drop in the
strength, to 309 MPa, and the DD11 weld strength
depletes even by 26 %, to 222 MPa, after the 480 h NSS
test. In contrast, no loss in the strength was found for
both aluminium alloys. Due to the severe corrosion
attack, the strength of the magnesium AZ31/AM50 weld
joint decreased significantly, by 78 %, down to 23 MPa,
Figure 18a.
FS test
The DD11-steel weld strength decreased to a value of
331 MPa (–6 %). However, for DD13 no loss of strength
was detected. Also, for the AW5083-O/AlMg5 and
AW6181-T4/AlSi5 weld joints no effect of corrosion
occurred after 480 h. The weld strength of AZ31
decreased sharply to 62 MPa (–40 %), Figure 18b.
3.4 Approximation of the base-material-corrosion
mass losses and rates
Weight measurements were performed on the 50 mm
× 100 mm base-material samples, see results in Figure
19. To calculate the mass loss per unit area (g/m2) For-
mula 1 was used. Therefore, the chemical composition of
the resulting corrosion products (MenXm) on the sample
surfaces must be known. Their respective compositions
from the literature are Fe2O3 for iron41, Al2O3 for
aluminum10,42 and Mg(OH)2 for magnesium8. According
to DIN 50905, part 243, the mass loss per unit area and
the corrosion rate can be calculated for the supposed
uniform surface corrosion. It must be noted that the alu-
minum samples showed irregular or localized corrosion.
For pitting corrosion, selective corrosion and intergra-
nular corrosion, DIN 50905, Part 344 describes the
evaluation of the corrosion samples by determining the
thinning. But in order to obtain comparability for all the
tested materials, a simplified uniform surface corrosion
was presumed. Applying Formula 2 and Formula 3, the
material-loss rates can be calculated43. The mass-loss
rate per unit area was calculated with difference ma from
mass loss ma at time tn reduced by the mass loss of the
previous evaluation time t(n – 1).
m
m
A
m
A
n A
m Aa
r, Me
r, X
=
−
=
+
⋅
⋅
⋅
Δ Δ
Formula 1: Mass loss per unit area 17
v
m
tdiff
ad
d
=
Formula 2: Mass-loss rate per unit area (g/m2h) 43
w
v
diff
diff= ⋅8 76.

Formula 3: Corrosion rate (mm/a) 43
Legend:
ma… Mass loss per unit area (g/m2)
–m… Mass loss (g)
+m… Mass increase (g)
A… Surface (m2)
n*Ar,Me… n · relative atomic mass of corrosion product
MenXm
m·Ar,X… m · relative atomic mass of non-metal X (oxy-
gen) in composition MenXm
vdiff… Mass loss rate per unit area (g/m2h)
dma… Differential mass loss per unit area (g/m2)
dt… Differential time between single tests (h)
… Base material density (g/cm3)
The calculation results are the following:
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Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 373
Figure 18: Welded-joint strength over the corrosion testing time; a)
NSS; b) FS test
Slika 18: Trdnost zvarjenega spoja v odvisnosti od ~asa korozijskega
preizku{anja; a) NSS-; b) FS-preizkus
3.4.1 NSS test
In Figure 19a the mass gains of all the examined
materials during the period of 480 h are shown. Four
samples were evaluated for each base material. Non-
grinded magnesium AZ31 (line 6) initially shows the
highest mass increase. After 240 h the mass decreases
due to the loss in the corrosion product. The mass
increase in the grinded AZ31 samples (line 5) is much
lower and even below the steel mass increase. The DD11
(line 3) and DD13 (line 4) steels show the same mass
gain up to 96 h. Thereafter, DD13 corrodes more signifi-
cantly. Aluminium AW5083-O (line 1) and AW6181-T4
(line 2) tend towards pitting and intergranular corrosion,
causing very low mass gains.
3.4.2 FS test
In Figure 19b the mass loss over the testing time for
all the tested materials is exhibited. It can be seen that
AZ31 (line 5) and DD13 (line 4) are forming adherent
corrosion layers up to 2 h. All the other materials
(AW6181-T4, AW5083-O and DD11) show neither an
increase nor a decrease. Up to 24 h all the samples gain
weight due to the formation of corrosion products.
Afterwards, until the test end, a steady weight loss for all
the samples is observed (except for the AW5083-O
increase up to 48 h). After 48 h steels DD11 (line 3) and
DD13 (line 4) sustain a strong weight loss, i.e., approxi-
mately 1.25 g. By comparison, the 0.3 g weight losses in
the aluminium alloys are quite low. The weight loss in
AZ31 after 480 h amounts to 0.45 g. While drying the
AZ31 samples, flaky corrosion products fall off the
samples that cannot be included in the mass loss.
3.4.3 Mass loss per unit area – the NSS test
The standardized mass losses after 240 h for steels
DD13 (1250 g/m2 – line 4), DD11 (1000 g/m2 – line 3)
and non-grinded magnesium AZ31 (800 g/m2 – line 6)
progress relatively linear, Figure 20a. At 480 h the
corrosion for DD11 decreases stronger than for DD13.
Non-grinded AZ31 shows a negative mass loss after 480
h due to the corrosion product falling off. The grinded
AZ31 (line 5) shows a constant and much lower mass
loss of just corrosion until the test end, whereat the
corrosion product adheres on the sample due to the low
surface roughness. AW5083-O (line 1) and AW6181-T4
(line 2) show a very low mass loss as expected.
3.4.4 Mass loss per unit area – the FS test
Figure 20b shows that steels DD11 (line 3) and
DD13 (line 4) have the highest mass loss. From 96 h
onwards the loss increases quite linearly up to
approximately 240 g/m2 after 480 h. For the materials
AW6181-T4 (25 g/m2), AW5083-O (50 g/m2) and AZ31
(87 g/m2) the mass loss is considerably lower.
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374 Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377
Figure 20: Mass loss per unit area over the 480 h test period: a) NSS
test, b) FS test
Slika 20: Zmanj{anje mase na enoto povr{ine med preizkusom 480 h;
a) NSS-preizkus, b) FS-preizkus
Figure 19: Base-material mass changes over the 480 h test period, a)
NSS test b) FS test
Slika 19: Spreminjanje mase osnovnega materiala med preizkusom
480 h, a) NSS-preizkus b) FS-preizkus
3.4.5 Corrosion rate – the NSS test
For the steels at the beginning of the NSS test, the
corrosion rates are quite high, i.e., around 12 mm per
year at instant 6 h. The non-grinded magnesium AZ31
exhibits a tremendous corrosion rate of 52 mm per year,
but the grinded AZ31 shows just 8.4 mm per year. To the
contrary, the aluminium alloys show low or immeasu-
rable corrosion rates (AW6181-T4: 2.2 mm per year;
AW5083-O: 0.0 mm per year ) at instant 6 h. By
continuing the test, the corrosion rates decrease more or
less to 240 h and fall down further at instant 480 h for all
the materials, Figure 21a. The reason for these
materials’ passivation is probably the adherent corrosion
products on the sample surfaces acting as corrosion
barriers. At the end of the NSS test, the steels (0.6 and 2
mm per year for DD11 and DD13) have the highest and
the aluminium alloys (AW5083-O and AW6181-T4, each
having 0.04 mm per year ) the lowest corrosion rate, see
Table 3.
3.4.6 Corrosion rate – the NSS test
The maximum corrosion rates were found at instant 6
h for AW6181-T4 (13 mm per year), AW5083-O (10 mm
per year), AZ31 (5.6 mm per year) and DD13 (0.9 mm
per year). Steel DD11 (1.5 mm per year) shows its maxi-
mum at instant 24 h. Thereafter, the corrosion rates of
the Al and Mg alloys fall significantly and only slightly
for the steels. At the end of the FS test AW6181-T4 (0.2
mm per year) and AW5083-O (0.3 mm per year ) show
the lowest corrosion rate. On the other hand, DD11 (0.5
mm per year ) and DD13 (0.7 mm per year) have a signi-
ficantly higher rate and AZ31 (1.4 mm per year) has the
highest rate, Figure 21b and Table 3.
4 DISCUSSION
The type of corrosion in the salt spray test (NSS test)
and fruit acid test (FS test) is as follows:
Aluminium alloys exhibit localized corrosion. It is
notable that the progress in pitting for AW5083-O and
intergranular corrosion for AW6181-T4 is reduced
during the testing time.
The grinded magnesium alloy AZ31 shows a stronger
tendency for localized corrosion in the FS test than in the
NSS test. The non-grinded AZ31 is not applicable as a
very heavy corrosion attack appears.
The corrosion behaviours of deep-drawing steels
DD11 and DD13 are quite similar. In the NSS test
mainly uniform corrosion or shallow pitting arises and in
the FS test a partial pitting corrosion is observed.
Note that the mass changes due to the corrosion in
the NSS test for the steels and magnesium alloy are two
to three times higher than in the FS test. Because of the
localized-corrosion behaviour of aluminium alloys their
mass change is marginal. Similarly, the corrosion rates at
certain testing times in the NSS test are multiple com-
pared to the FS test for the steels, Table 3.
It is remarkable that in the NSS test adherent corro-
sion layers are formed on all the materials and the sam-
ples gain in weight. In contrast, the corrosion products
drain off in the FS test making the samples lose weight.
The corrosion behaviour of the welded joint
AW5083-O in the NSS and FS tests is similar. Here the
pitting and localized corrosion was found at the fusion
line and also at the weld edges on the AlMg5 weld
metal. Amazingly, the AW6181-T4 welded joints
showed no corrosion at the fusion line or on the AlSi5
weld metal, but just the IG corrosion on the base
material.
The corrosion attack on the AZ31 magnesium welded
joint was stronger in the NSS test than in the FS. In the
former a strong and partially localized corrosion on the
AM50 weld metal and on the base material occurred. In
the latter only shallow pitting corrosion was found.
The welded steel joints showed an intense corrosion
attack at the fusion line in the NSS and FS tests. The
corrosion of the G3Si1 weld metal was selective in NSS
and took the form of shallow pitting corrosion in the FS
test.
The degradation of the tensile-shear strength of the
steel welding joints is stronger in the NSS (up to 25 %)
than in the FS test. Due to a very high strength loss of
the AZ31 magnesium this joint is not applicable. It is a
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Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 375
Figure 21: Corrosion-rate behaviour over the 480 h test period: a)
NSS test b) FS test
Slika 21: Spreminjanje hitrosti korozije med preizkusom 480 h: a)
NSS-preizkus, b) FS-preizkus
benefit of aluminium alloys that they show no influence
of corrosion on the welding-joint strength.
5 CONCLUSIONS
Basically, the results of this investigation program
allow a technical comparison between the influences of
the two relevant corrosion media on the household cool-
ing compressor shells.
The currently used deep-drawing steels cannot be
used without an anti-corrosion coating (KTL). Due to a
low strength and even a poor corrosion resistance, it is
not advisable to use AZ31 as a housing material for
household cooling compressors.
The overall result regarding weldability, corrosion
resistance and joint strength proves that aluminium
alloys can bring improvements to the future compressor
shells.
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376 Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377
Table 3: Comparison of the mass-loss rates v/(g/(m2 h)) and corrosion rates w/mm per year in the NSS and FS test intervals
Tabela 3: Primerjava hitrosti izgube mase v/(g/(m2 h)) in korozijska hitrost w/mm na leto pri NSS- in FS-intervalih preizku{anja
Material
Neutral salt spray test (NSS test) Fruit acid spray test (FS test)
96–168 168–240 240–480 96–168 168–240 240–480
v
g/(m2 h)
w
(mm/y)
v
g/(m2 h)
w
(mm/y)
v
g/(m2 h)
w
(mm/y)
v
g/(m2 h)
w
(mm/y)
v
g/(m2 h)
w
(mm/y)
v
g/(m2 h)
w
(mm/y)
DD11 1.80 2.01 5.53 6.16 0.54 0.60 0.84 0.94 0.39 0.44 0.47 0.52
DD13 2.71 3.02 8.01 8.91 1.85 2.06 0.62 0.69 0.29 0.33 0.58 0.65
AW5083-O 0 0 0.02 0.05 0.01 0.04 0.23 0.76 0.02 0.07 0.1 0.33
AW6181-T4 0.01 0.02 0.02 0.07 0.01 0.04 0.15 0.48 0.07 0.24 0.07 0.22
AZ31 0.14 0.73 0.24 1.22 0.27 1.35 0.13 0.66 0.17 0.84 0.27 1.37
35 G. Svenningsen et al., Effect of low copper content and heat treat-
ment on intergranular corrosion of model AlMgSi alloys, Corrosion
Science, 46 (2006), 226–242
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tions, Second English Edition. s.l., Cebelcor Verlag, 1974, 314
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Berlin 1969
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gen-Teil 2, Beuth Verlag GmbH, Berlin 1987
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– Teil 3, Beuth Verlag GmbH, Berlin 1987
K. G. KERSCHBAUMER et al.: WELDED ALUMINIUM AND MAGNESIUM ALLOYS – CORROSION AND MECHANICAL ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 363–377 377

J. VASILJEVI] et al.: WATER-VAPOUR PLASMA TREATMENT OF COTTON AND POLYESTER FIBRES
WATER-VAPOUR PLASMA TREATMENT OF COTTON
AND POLYESTER FIBRES
OBDELAVA BOMBA@NIH IN POLIESTRSKIH VLAKEN
S PLAZMO VODNE PARE
Jelena Vasiljevi}1, Marija Gorjanc1, Rok Zaplotnik2, Alenka Vesel2,
Miran Mozeti~2, Barbara Simon~i~1
1Department of Textiles, Faculty of Natural Sciences and Engineering, University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia
2Department of Surface Engineering and Optoelectronics, Jo`ef Stefan Institute, Jamova cesta 39, 1000 Ljubljana, Slovenia
marija.gorjanc@ntf.uni-lj.si, barbara.simoncic@ntf.uni-lj.si
Prejem rokopisa – received: 2012-10-05; sprejem za objavo – accepted for publication: 2012-11-13
This study deals with the treatment of cotton and polyester fibres with a low-pressure, inductively coupled RF plasma, in which
the water vapour from the fibres was used as a plasma-forming gas. Cotton, which is a hydrophilic, natural, cellulose fibre and
polyester, which is a hydrophobic, synthetic, polyethylene terephthalate fibre, were used. Plasma characteristics during the
treatment were investigated using optical emission spectroscopy (OES). The morphological and chemical changes in the fibre
surfaces induced by plasma treatment were analysed using atomic force microscopy (AFM), scanning electron microscopy
(SEM) and X-ray photoelectron spectroscopy (XPS). The optical emission spectra showed the presence of OH and H radicals at
the beginning of the plasma treatment, whereas a CO Angstrom band appeared in the spectra recorded during the plasma
treatment of both fibres. The cotton fibre roughness showed a three-fold increase after the plasma treatment that increased the
surface area by approximately 8 %. The changes in the polyester fibre roughness were much less distinct and the surface area
increased by approximately 3 %. The plasma treatment induced an increase in the O/C atomic ratio by approximately 43 % for
cotton and 56 % for polyester fibres. The etching action of the water-vapour plasma was thus found to be more effective on the
surface of cotton fibres than on polyester fibres. However, the water content of polyester fibres was sufficiently high to cause an
oxidation of the fibre surface that was even higher than the oxidation of cotton.
Keywords: water-vapour plasma, effectiveness of treatment, cotton, polyester, water content, morphological and chemical
changes
Raziskava vklju~uje obdelavo bomba`nih in poliestrskih vlaken z nizkotla~no induktivno sklopljeno RF-plazmo vodne pare.
Izvir vodne pare kot delovnega plina so bila vlakna bomba`a, predstavnika hidrofilnih naravnih celuloznih vlaken, in vlakna
poliestra, predstavnika hidrofobnih sinteti~nih polietilenteraftalnih vlaken. Lastnosti plazme med obdelavo tekstilnih vzorcev so
bile preiskane z opti~no emisijsko spektroskopijo (OES). Morfolo{ke in kemijske spremembe povr{in plazemsko obdelanih
vlaken so bile analizirane z mikroskopijo na atomsko silo (AFM), z vrsti~no elektronsko mikroskopijo (SEM) in rentgensko
fotoelektronsko spektroskopijo (XPS). Opti~ni emisijski spektri so na za~etku obdelave s plazmo pokazali prisotnost OH in H
radikalov. Med plazemsko obdelavo obeh vrst vlaken so bili v opti~nih emisijskih spektrih opazni tudi trakovi, ki izvirajo iz
prehodov radikalov CO. Po obdelavi s plazmo se je hrapavost povr{ine bomba`nega vlakna trikrat pove~ala, kar je vplivalo na
pove~anje specifi~ne povr{ine za pribli`no 8 %. Spremembe v hrapavosti so bile pri poliestrskem vlaknu manj izrazite,
specifi~na povr{ina vlaken se je pove~ala za pribli`no 3 %. Po obdelavi s plazmo se je na povr{ini obeh vrst vlaken pove~alo
razmerje O/C, in sicer za 43 % pri bomba`u in za 56 % pri poliestru. Glede na vsebnost vode v tekstilnih vzorcih je imela
obdelava s plazmo vodne pare ve~ji u~inek jedkanja na bomba`nih kot na poliestrskih vlaknih. Kljub temu je poliester vseboval
dovolj vodne pare, da se je povr{ina vlaken v plazmi oksidirala. Oksidacija na poliestru je bila celo ve~ja kot na bomba`u.
Klju~ne besede: plazma vodne pare, u~inkovitost obdelave, bomba`, poliester, vsebnost vode, morfolo{ke in kemijske
spremembe
1 INTRODUCTION
Non-equilibrium gaseous plasma is a unique techno-
logy for treating the surfaces of fibrous polymers without
affecting their bulk properties.1–9 Accordingly, plasma
systems using different power sources, pressures, elec-
trode configurations and gasses have been used to
generate plasma discharges and treat polymers.10–20
Oxygen-containing plasma increases the specific surface
area and surface energy of polymers, which results in an
incorporation of oxygen-containing groups, such as
C–O, O–C=O and C=O.21–26 Plasma discharges in water
vapour can potentially be used in a wide range of
applications.27–32 Water-vapour plasma is superior to the
other plasma-forming gasses because of its unique pro-
perties, i.e., an extremely high enthalpy, environmentally
safe conditions, a relatively low cost and an endless
amount of plasma-forming gas.33 Water-vapour plasma
generates a high concentration of OH radicals that can
further dissociate to H and O radicals,34–36 although the
probability of a dissociation of OH radicals is lower than
that of a dissociation of water molecules (some of the
OH radicals remain undissociated).36 The first effect of a
plasma treatment is the functionalization of the polymer
surface, which is followed by etching reactions. These
reactions are initiated by an H-atom abstraction and
formation of a free radical.35,37–40
Cotton and polyester are the two most important and
widely used polymers in the textile industry. They differ
from each other in their chemical and morphological
structures, and the amounts of water content. There are
several publications dealing with the influence of diffe-
Materiali in tehnologije / Materials and technology 47 (2013) 3, 379–384 379
UDK 677:677.21:677.494.674 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(3)379(2013)
rent plasma discharges on the surface properties of
cotton or polyester, but these two polymers were rarely
investigated together under the same plasma parameters.
The aim of this study was to investigate the surface
changes of cotton and polyester after a treatment with
water-vapour plasma under the same plasma parameters.
The water vapour from the fibres was used as a plasma-
forming gas. The low-pressure plasma was chosen as an
environmentally friendly (pre)treatment of the textiles,
water vapour was chosen as an environmentally safe and
cheap plasma-forming gas that is already present in the
fibres, and the same plasma parameters were chosen to
investigate the surface changes occurring on the two
different yet most commonly used polymers. Additio-
nally, plasma treatment was observed using OES, where
the optical spectra were recorded during the treatment of
textiles. The changes to the fibre surfaces were investi-
gated using SEM, AFM and XPS analyses.
2 EXPERIMENTAL WORK
2.1 Materials and methods
Scoured, bleached and mercerised cotton woven
fabric (119 g/m2) made by Tekstina, d. d., Ajdov{~ina,
and washed polyester woven fabric (67 g/m2) made by
Velana, d. d., Ljubljana, were used in the study.
Low-pressure, inductively coupled, radiofrequency
(RF) plasma was used for the treatment. Our RF reactor
scheme was previously presented in detail.27 The plasma
was created using a fixed frequency of 27.12 MHz, an
output power of approximately 5 kW, a pressure of 20 Pa
and a current of 0.4 A. Water vapour was used as a work-
ing gas. The fabric was the source of the water vapour.
The fabric samples were treated for 30 s.
2.2 Analyses and measurements
The water content of the fabric samples was deter-
mined according to the standard SIST ISO 6741-1:1996.
The plasma treatments were observed using optical
emission spectrometry. An Avantes AvaSpec-3648 opti-
cal spectrometer with a 3648-pixel CCD detector array
and a 75 cm focal length was used. This spectrometer
records the optical emission spectra in the wavelength
range of 200–1100 nm with a resolution of 0.5 nm.
The sample morphology was evaluated using a JEOL
JSM 6060 LV scanning electron microscope. The sam-
ples were coated with a thin layer of gold before obser-
vation.
The sample surface topography was evaluated using
an atomic force microscope (AFM/MFM – Veeco
Dimension 3100) in the contact mode. Surface scans of
1 μm2 square areas were performed at 22 °C in atmo-
sphere for different sample positions using a scanning
rate of 1.51 Hz. Images with a resolution of 256 × 256
lines were obtained using the Nanoscope software with a
Flatten filter. From the AFM analyses, the mean rough-
ness (Ra; the arithmetic average of the deviation from the
centre plane), root-mean-square roughness (Rms; the
standard deviation of the Z-value within a given area)
and surface area (SA; a three-dimensional given region)
were calculated as the mean values of 10 AFM scans of
different sample regions. The standard error of the mean
was calculated to perform statistical analyses.
The sample-surface chemical compositions were
analysed with X-ray photoelectron spectroscopy using a
PHI-TFA XPS spectrometer (Physical Electronics Inc).
The analysed areas had a diameter of 0.4 mm and a
depth of approximately 3–5 nm. The sample surfaces
were excited with the X-ray radiation from a mono-
chromatic Al K
 source at a photon energy of 1486.7 eV.
The sample-surface chemical compositions were quanti-
fied on the basis of the XPS peak intensities measured
for two different spots on the sample using the MultiPak
v7.3.1 software from Physical Electronics that was
supplied with the spectrophotometer.
3 RESULTS AND DISCUSSION
Inductively coupled RF plasma excited at the ulti-
mate pressure emitted the spectra shown in Figure 1. As
water remained in the samples, the residual atmosphere
was mostly water vapour. Therefore, hydrogen emission
lines were the main lines that were visible in the optical
emission spectra recorded at the beginning of the treat-
ment (Figure 1).
Hydrogen is a good emitter of radiation, which is
why the intensities of the Balmer series emission lines
are so high. Water molecules dissociate into the hydro-
gen and oxygen atoms, but in the OES spectra of the
inductively coupled plasma at the ultimate pressure, the
oxygen lines are not present, whereas the excitation ener-
gy of oxygen atoms is much higher than the excitation
energy of hydrogen atoms. Nitrogen emission bands are
J. VASILJEVI] et al.: WATER-VAPOUR PLASMA TREATMENT OF COTTON AND POLYESTER FIBRES
380 Materiali in tehnologije / Materials and technology 47 (2013) 3, 379–384
Figure 1: OES spectra of the plasma, generated at the ultimate
pressure before etching
Slika 1: OES-spektri plazme, ustvarjene pri kon~nem tlaku pred jed-
kanjem
present in the spectra because of the small leakages in
the system.
The optical spectra recorded during the treatment of
the samples (Figures 2 and 3) revealed several CO Ang-
strom-band emission lines as well as a broad continuum
between 400 nm and 700 nm that most likely resulted
from a partial overlapping of the radiative transitions
within CO molecules. The CO emission lines are attri-
buted to the etching of the cotton and polyester samples.
OES is a qualitative technique and, therefore, the density
of particles cannot be determined from these measure-
ments. However, CO is a poor emitter and from the high
CO emission lines, we can conclude that the etching was
efficient.
The time evolution of the etching was recorded after
the CO Angstrom band (0, 2) emission peak was obser-
ved. To eliminate the effects of different spectrometer
optical-fibre positions, the CO emission line (519 nm)
was normalised with the H line. The time evolution of
the normalised CO emission line during the sample treat-
ment is presented in Figure 4.
Better etching was achieved for the cotton samples
when compared with the polyester samples. The etching
rate for both samples increased monotonically with time.
This etching increase was attributed to the thermal
effects: the samples underwent heating during the treat-
ment that increased the etching rate.
The etching effects of the water-vapour plasma treat-
ment were further examined using SEM and AFM
analyses. The SEM images of cotton and polyester fibres
before and after the plasma treatment are presented in
Figure 5. Under the used plasma parameters, the impuri-
ties on the cotton surface were cleaned and removed
(Figures 5a and c). The same plasma parameters did not
induce such changes to the polyester surface (Figures 6a
and c). The plasma treatment did not cause any visible
J. VASILJEVI] et al.: WATER-VAPOUR PLASMA TREATMENT OF COTTON AND POLYESTER FIBRES
Materiali in tehnologije / Materials and technology 47 (2013) 3, 379–384 381
Figure 5: SEM images of: a), b) untreated cotton and c), d) plasma-
treated cotton fibres; 500-times magnification was used for a) and c),
6500-times magnification for b) and d)
Slika 5: SEM-posnetki vlaken: a), b) neobdelanega bomba`a ter c), d)
s plazmo obdelanega bomba`a; a) in c) 500-kratna pove~ava in b) in d)
6500-kratna pove~ava
Figure 3: OES spectra of the plasma during the etching of polyester
Slika 3: OES-spektri plazme med jedkanjem poliestrskega vzorca
Figure 4: Time evolution of the CO emission peak (519 nm),
normalised with the H line during the treatment of: a) cotton and b)
polyester
Slika 4: ^asovni razvoj emisije vrha CO (519 nm), normaliziranega s
H-~rto, med obdelavo vzorca: a) bomba`, b) poliester
Figure 2: OES spectra of the plasma during the etching of cotton
Slika 2: OES-spektri plazme med jedkanjem bomba`nega vzorca
morphological surface changes on cotton (Figures 5b
and d) or on polyester fibres (Figures 6b and d), indi-
cating that the bulk properties of both types of fibres
remained unchanged.
AFM analyses showed the changes in the nanotopo-
graphy that were induced by the water-vapour plasma
treatment of both samples (Figure 7).
The quantitative evaluation of the nanotopography
changes is presented in Figure 8 in terms of Ra and SA.
As the calculated Rms values had the same trend as the Ra
values, they are not presented in the paper. The plasma
treatment produced a three-fold Ra increase in the cotton
fibres, but did not cause significant changes in the poly-
ester fibres (Figure 8a). The calculated SA values ob-
tained after the plasma treatment were also higher for
cotton than for polyester fibres (Figure 8b). For cotton
and polyester, the SA values increased by approximately
8 % and 3 %, respectively. These results are in agree-
ment with the obtained OES results. A higher etching
effectiveness of the water-vapour plasma during the
cotton treatment was caused by a different water content
of the sample, i.e., 6.9 % for cotton and 0.5 % for poly-
ester. A higher fibre water content leads to a higher level
of reactive plasma species in the discharge, which
contributes to an enhanced etching effect.
The XPS analyses (Table 1) suggest that the oxida-
tion of plasma-treated cotton and polyester is not depen-
dant on their water contents. After a plasma treatment,
the atomic concentration of carbon (C1s, 285 eV)
decreased and the atomic concentration of oxygen
(O1s, 533 eV) increased in both cases. The calculated
J. VASILJEVI] et al.: WATER-VAPOUR PLASMA TREATMENT OF COTTON AND POLYESTER FIBRES
382 Materiali in tehnologije / Materials and technology 47 (2013) 3, 379–384
Figure 7: AFM scans of: a) untreated cotton, b) plasma-treated cotton,
c) untreated polyester and d) plasma-treated polyester fibres
Slika 7: AFM-posnetki vlaken: a) neobdelanega bomba`a, b) s plazmo
obdelanega bomba`a, c) neobdelanega poliestra in d) s plazmo
obdelanega poliestra
Figure 8: a) Mean roughness, Ra, and b) surface area, SA, of cotton
and polyester samples
Slika 8: a) Srednja vrednost hrapavosti, Ra, in b) specifi~na povr{ina,
SA, bomba`nih in poliestrskih vzorcev
Table 1: Elemental compositions of the fabric-sample surfaces deter-
mined with a XPS analysis
Tabela 1: Elementarna sestava povr{ine vzorcev tkanin, dobljena z
XPS-analizo
Sample
Elemental composition
in mole fractions (%) Atomic
ratio O/C
x(C) x(O)
Untreated cotton 69.4 30.6 0.44
Plasma treated cotton 61.3 38.7 0.63
Untreated polyester 72.0 28.0 0.39
Plasma treated polyester 62.2 37.8 0.61
Figure 6: SEM images of: a), b) untreated polyester and c), d) plas-
ma-treated polyester fibres; 500-times magnification was used for a)
and c), 6500-times magnification for b) and d)
Slika 6: SEM-posnetki vlaken: a), b) neobdelanega poliestra ter c), d)
s plazmo obdelanega poliestra; a) in c) 500-kratna pove~ava in b) in d)
6500-kratna pove~ava
increases of the O/C atomic ratios were 43 % for cotton
and 56 % for polyester. Despite a low water content of
the polyester sample, the fibre oxidation during the
plasma treatment was successful.
4 CONCLUSIONS
A low-pressure water-vapour plasma was used for
treating cotton and polyester fabrics. The source of the
water vapour was the fabric itself. The basis of this
research was to investigate the surface changes in the
two chemically and morphologically different textiles
when the same plasma parameters were used. Before and
after the plasma treatment the surface properties of
cotton and polyester fibres were evaluated and compared
using the XPS, SEM and AFM analyses. The plasma
treatment of the textiles was observed using an OES
analysis. The results showed that the etching rate
increases with the treatment time for both types of
polymers due to the sample heating during the plasma
treatment. The etching effect of the low-pressure water-
vapour plasma was more pronounced for the hydrophilic
cotton fibres than for hydrophobic polyester fibres. The
cleaning effect on the surface of the cotton fabric was
observed, when the surface impurities were removed.
Under the used plasma parameters, surface morphology
of the fibres remained unchanged for both types of
polymers, indicating that the bulk properties of the fibres
remained undamaged. The water content of both types of
polymers was sufficiently high to achieve a good
oxidation of the fibre surfaces.
This preliminary research is very important for
understanding what is happening during plasma treat-
ment of cotton and polyester and how the same plasma
parameters influence the surface changes made to these
two most important and widely used polymers. Since
cotton and polyester are also used together as a cotton/
polyester blend, the results will help us with our further
research focusing on the water-vapour plasma (pre)treat-
ment of blends.
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384 Materiali in tehnologije / Materials and technology 47 (2013) 3, 379–384
A. SALIMI et al.: INFLUENCE OF SAMPLE DIRECTION ON THE IMPACT TOUGHNESS ...
INFLUENCE OF SAMPLE DIRECTION ON THE IMPACT
TOUGHNESS OF THE API-X42 MICROALLOYED STEEL
WITH A BANDED STRUCTURE
VPLIV USMERJENOSTI VZORCEV NA UDARNO @ILAVOST
MIKROLEGIRANEGA JEKLA API-X42 S TRAKAVO
MIKROSTRUKTURO
Ahmadreza Salimi1, Hossein Monajati Zadeh1, Mohammad Reza Toroghinejad2,
Davod Asefi1, Amir Ansaripour1
1Department of Materials Engineering, Najafabad Branch, Islamic Azad University, P.O. Box 517, Isfahan, Iran
2Department of Materials Engineering, Isfahan University of Technology, Isfahan, Iran
salimy.ahmadreza@gmail.com
Prejem rokopisa – received: 2012-08-13; sprejem za objavo – accepted for publication: 2012-10-23
The layering of a microstructure parallel to the direction of the material flow during the hot working process is called banding.
In the present paper, the severity of ferrite-pearlite banding in the API-X42 microalloyed steel and its effects on the impact
energy are studied. Specifically, the impact toughness is examined in the cases of banded and non-banded samples along two
directions, perpendicular and parallel to the rolling, and the obtained results are compared. Metallographic examinations,
together with the Charpy impact tests at 0 °C and –18 °C, were done for both directions, parallel and perpendicular to the
rolling. The results showed a dependence of the impact energy on the sample direction relative to the rolling to be larger with
the banded than the non-banded sample. The difference between the impact energies for the directions parallel and
perpendicular to the rolling was also noticed to be caused by the increasing anisotropy index.
Keywords: API-X42 steel, pearlite-ferrite banding, impact energy
Usmerjanje mikrostrukture vzporedno s smerjo vro~ega preoblikovanja se imenuje trakavost. V ~lanku je predstavljena {tudija
preprostosti feritno-perlitnih pasov v mikrolegiranem jeklu API-X42 in njihov vpliv na udarno energijo. Specifi~no je
primerjana udarna `ilavost med vzorci s trakavo in netrakavo mikrostrukturo vzdol` dveh smeri, pravokotno in vzporedno s
smerjo valjanja. Opravljene so bile metalografske preiskave in udarni preskusi Charpy pri 0 °C in –18 °C vzporedno in
pravokotno na smer valjanja. Rezultati so pokazali odvisnost udarne energije glede na usmerjenost vzorca pri valjanju, relativno
bolj v trakavih kot v netrakavih vzorcih. Pri nara{~ajo~em indeksu anizotropije je bila opa`ena razlika med energijo udarca pri
vzporedni in pravokotni smeri valjanja.
Klju~ne besede: jeklo API-X42, perlitno-feritni trakovi, udarna energija
1 INTRODUCTION
The layering of a microstructure parallel to the
direction of the material flow during the hot working
process is called banding. Generally, banding is
classified into two major categories. The microscopic
bands include deformation bands, transformation bands
and shear bands, and the macroscopic bands include
carbide banding in tool steels, layered ferrite-pearlite
structure of rolling in low-carbon alloy steels and
martensite banding in heat-treated alloy steels.1–4
Because of these various micro and macro features, there
is no universally accepted definition of banding. More-
over, there are clearly various mechanisms that can cause
these structures.
Ferrite-pearlite banding may occur due to a segre-
gation of some alloying elements during solidification
after casting and hot-working processes.5 When steel is
slowly cooled from the austenite region, the pro-eutec-
toid ferrite is formed initially in the areas with a rela-
tively low number of austenite-stabilizing elements,
whereas pearlite is formed in the areas with more
austenite-stabilizing elements after being cooled down to
the temperatures below the eutectoid line, creating a
banded microstructure containing successive pearlite and
ferrite areas.6
There have been several studies on the effect of
banding on mechanical properties7–9. Working on heavily
banded 0.3 % carbon steel, Jatczak et al.10 found little or
no effect on the anisotropy of tensile properties, while a
significant anisotropy of the reduction in area and impact
properties was discovered. They also observed a very
small change in mechanical properties in longitudinal
direction as well as in impact properties and ductility in
transverse direction due to homogenization.
Grange11 found that both banding microstructure and
longitudinally directed inclusions cause anisotropy in the
mechanical properties of 0.025 % C and 1.5 % Mn steel,
eliminating the decrease in anisotropy caused by
banding; however, this decrease is trivial if numerous
inclusions are elongated in longitudinal direction. In
some studies, the effect of banding and specimen orien-
tations on the fracture toughness has been investigated
and it has been shown that banding has a significant
Materiali in tehnologije / Materials and technology 47 (2013) 3, 385–389 385
UDK 539.42:621.77.014.2 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(3)385(2013)
effect on the rolling-plane anisotropy.12–14 However, little
has been done to examine the differences between
banded and non-banded samples in different directions.
In the previous researches, the samples with a banded
structure were selected, and after the study of structural
and mechanical properties, the same steel was heat
treated, by normalizing or annealing it, to remove micro-
structural banding.2–15 The potential problem of this
method is that this treatment makes it possible to prepare
samples of the same steel in a virtually non-banded ver-
sus severely banded conditions; however, the differences
in the chemical-composition distribution, grain size and
inclusion morphology during heat treatment are inevi-
table. So, they can affect the accuracy of the results. In
this study, the impact properties of several API-X42 steel
samples with different bandy degrees of the ferritic-
pearlite structure after hot rolling have been investigated.
No heat treatment was done on the samples to reduce the
banding phenomenon.
2 EXPERIMENTAL PROCEDURE
In this study, the initial production data of 50 samples
of the API-X42 steel were obtained. Among them, 16
samples with the same chemical composition were
selected. The chemical composition in mass fractions
(%) is 0.12 C, 0.905 Mn, 0.21 Si, 0.007 P, 0.003 S and
25 N (μg/g). The samples were investigated with optical
metallography.
Metallographic specimens were prepared in accor-
dance with the guidelines and recommended practices
given by ASTM-E3 Methods. Their images provided by
an optical microscope at the magnification of 100 and
500 were also taken.
To investigate just the banding effects and remove the
other effects of metallurgical variables, samples with the
same grain size and chemical composition and with very
low amounts of inclusions have been selected.
The banding in API-X42 is the ferrite-pearlite
banding. Hence, the banded samples and non-banded
ones were separated from each other by using the
Assessing the Degree of Banding or Orientation of
Microstructures standard (ASTM-E1268). The aniso-
tropy index (AI) was estimated from the following
equation:
AI = NL/NLII (1)
where NL and NLII are the mean numbers of the feature
interceptions with the test lines respectively perpen-
dicular and parallel to the deformation direction per
length unit of the test lines. For a randomly oriented,
non-banded microstructure, AI has a value of one. As
the degree of orientation or banding increases, AI
increases, too.
For the impact testing (according to the ASTM-E23
standard), three samples from each plate in a direction
perpendicular to rolling and three samples in parallel
with the rolling direction were prepared as shown in
Figure 1. (At the state A, the test piece is perpendicular
to the rolling direction and the notch is parallel to
rolling. At the state B, the test piece is parallel to the
rolling direction and the notch is perpendicular to rolling
and the pendulum strikes the test piece in the direction
parallel to rolling).
The Charpy impact tests at 0 and –18 °C were done.
The amount of the absorbed energy was determined for
each test piece. Finally, the mean of the three results for
each temperature was reported as the final result. The
fracture surface was coated with nickel to be prepared
for examining the crack-propagation path in normal view
of the fractured face. Nickel prevents damaging the
studied fracture surface. Then the back of the broken
test-piece notch was investigated with SEM microscopy.
3 RESULTS AND DISCUSSION
Figure 2 shows two metallographic images used for
determining AI for the two cases of highly and poorly
banded microstructures. The results for the classified
specimens after the metallography, the anisotropy index
and the impact energy, are shown in Table 1.
Table 1: Relationship between the impact energy and the anisotropy
index for the two temperatures of 0 °C and –18 °C
Tabela 1: Odvisnost med udarno energijo in indeksom anizotropije za
dve temperaturi, 0 °C in –18 °C
No. of piece AI
Impact energy, J
(–18°C) (0°C)
1A 2.07 46.4 63.0
2A 2.18 55.4 67.5
3A 1.57 67.9 72.3
4A 1.55 67.4 74.0
1B 2.07 84.0 87.9
2B 2.18 88.8 89.7
3B 1.57 78.6 83.0
4B 1.55 79.6 85.0
In accordance with the banding standard, the optimal
state is achieved when AI is equal to 1. It means that the
A. SALIMI et al.: INFLUENCE OF SAMPLE DIRECTION ON THE IMPACT TOUGHNESS ...
386 Materiali in tehnologije / Materials and technology 47 (2013) 3, 385–389
Figure 1: Schematic representation of the sample direction relative to
the plate rolling direction
Slika 1: Shematski prikaz smeri vzor~enja in usmerjenost plo{~e
glede na smer valjanja
number of NL must be equal to the number of NLII. By
increasing it to a number higher than 1, NL is higher
than NLII, and the microstructure becomes more banded.
According to Table 1, all the measured values of AI
are greater than 1. This indicates that all the samples are
banded. But two series of more banded specimens
containing the test pieces 1A, 2A, 1B, 2B and two series
of less banded specimens containing the test pieces 3A,
4A, 3B, 4B can be singled out.
Three-dimensional metallographic microstructure
images of samples 4A and 2B are shown in Figures 3a
and 3b, respectively, as the samples with a poorly and a
highly banded microstructures.
As can be seen, the banding phenomenon is more
visible in the direction parallel to the rolling cross-
section. The relationship between the impact energy and
AI for the eight test pieces examined in the directions of
A and B, is shown in Figure 4.
According to Figure 4, in the samples of series A, the
impact energy is reduced with an increase in AI. As the
samples of series B behave differently, an increase in AI
causes an increase in the impact energy. About a 40 %
increase in AI results in about a 10 % increase in the
impact energy of B samples and about a 25 % decrease
in A samples.
This increasing and decreasing of the impact energy
caused by the increasing banding can be attributed to the
crack-growth path in the banding layers of the impact
sample. Schematic representations of the crack-growth
paths from samples A and B are shown in Figure 5.
According to this figure, in the A sample the crack
A. SALIMI et al.: INFLUENCE OF SAMPLE DIRECTION ON THE IMPACT TOUGHNESS ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 385–389 387
Figure 3: Three-dimensional images of metallographic samples 4A
and 2B as the two samples with a low and a high AI
Slika 3: Tridimenzionalna slika metalografskega vzorca 4A in 2B kot
dva vzorca z nizkim in visokim indeksom anizotropije AI
Figure 2: Two examples of a: a) banded structure and b) non-banded
structure, for which the anisotropy index was determined
Slika 2: Primer: a) trakave mikrostrukture in b) mikrostrukture brez
trakavosti, za katero je bil dolo~en indeks anizotropije
Figure 4: Dependence of the impact energy on the sample direction
(A and B according to Figure 1) and AI at –18 °C
Slika 4: Odvisnost udarne energije od usmerjenosti vzorca (A in B na
sliki 1) in indeksa anizotropije AI pri –18 °C
moves along the path of the layers placed on each other,
while in the B sample the crack movement is perpen-
dicular to the layers and is encountered by different
layers on its path. So, a stronger banding is more harmful
for the A sample and more useful for the B sample. This
is because more layers on the crack-growth path can
absorb more energy.
According to Table 1 and Figure 3, the 1B and 2B
test pieces that have a high AI, absorbing more energy
before the fracture, while the 3B and 4B test pieces that
have a low AI absorbing less energy and allowing the
fracture to occur. In addition, the minimum value is
related to the 3B sample having the lowest AI. It means
that banding can be useful to impact properties if the
pendulum impact on the test piece is in the direction
perpendicular to the rolling. A crack can deviate from its
main path due to grain boundaries, flow lines, inclusions
and banding. It can be seen in Figure 2 that in the
banded samples the distance between the ferrite and
pearlite is smaller than in the non-banded samples and
the crack is forced to encounter many more phases on its
path. Also, the phase continuity is a more important
factor in the case of highly banded samples than in the
case of poorly banded samples. The reason for this is the
fact that when phases are continuous, a crack is forced to
encounter many more phases on its path. On the other
hand, due to the discontinuity of the phases in the poorly
banded samples, a crack may follow a longer path before
encountering another phase. This can be better seen in
Figure 6, which shows a SEM image of the section
perpendicular to the fracture surface of sample 2. This
figure implies that the crack path has been serrated while
crossing the ferrite-pearlite banded structure.
The crack path on this figure is perpendicular to the
plate. So, during the crossing, it has been involved with
the structures of ferrite and pearlite. This caused a
change in the path and can be the reason for an increase
in the absorbed impact energy in B similar to the banded
samples.
The A banding structure can also be compared to a
fiber composite or a composite with constant reinforcing
particles, in which the reinforcing phase is located in the
field phase. The fracture in the direction perpendicular to
A. SALIMI et al.: INFLUENCE OF SAMPLE DIRECTION ON THE IMPACT TOUGHNESS ...
388 Materiali in tehnologije / Materials and technology 47 (2013) 3, 385–389
Figure 7: Difference in the impact energies for the A and B samples at
–18 °C and 0 °C
Slika 7: Razlika v udarni energiji pri dveh vzorcih A in B pri –18 °C
in 0 °C
Figure 5: Schematic image representing the crack-growth path in the
A and B banded samples
Slika 5: Shematski prikaz poti rasti razpoke v trakavih vzorcih A in B
Figure 6: SEM image of a crack-growth path in the cross-section
perpendicular to the fracture surface of sample 2B
Slika 6: SEM-posnetek poti rasti razpoke na preseku pravokotno na
povr{ino preloma vzorca 2B
the fibers is more serious than the one in the parallel
direction. Figure 7 shows the differences in the impact
energies for the two directions, perpendicular and parallel,
of the rolling.
In a study on an aluminum composite, in which the
SiC reinforcing particles are used in two forms, an
elongated and a random orientation in the aluminum
matrix, the sample with a random distribution showed no
difference in the fatigue crack growth in the two
directions, the perpendicular and parallel, while the
sample with an elongated reinforcement showed an
obvious difference.16
Also, the repeated impact tests at the zero tempera-
ture indicated a similar behavior shown in this diagram,
and the diagram in Figure 4 shows that with the
increasing AI the differences in the impact energies for
the two directions, the perpendicular and parallel,
increase from 11 J to 37 J. Therefore, the direction of the
samples has a significant effect on the banded samples,
reflected in the impact-test results. Due to a more
homogeneous microstructure of the non-banded samples
or the samples with a lower AI, the mechanical pro-
perties are not very dissimilar for different directions and
are close to the homogeneous state.
However, at the temperature of –18 °C the amount of
the impact energy in the banded samples was slightly
smaller than at the temperature of 0 °C. This shows that
banding has a larger effect on the impact properties at
lower temperatures.
4 CONCLUSION
1. Impact energy shows an obvious dependence on the
sample direction. As AI increases (which is an index
of banding), the impact energy decreases in the
rolling direction, while it increases in the direction
perpendicular to rolling.
2. With the increasing AI the differences in the impact
energies for the two directions, perpendicular and
parallel, increase.
3. At lower temperatures banding has a greater influ-
ence on impact properties.
Acknowledgments
The technical and financial support of the Mobarkeh
Steel Complex to the present research is gratefully
acknowledged.
5 REFERENCES
1 ASTM International, Metals-Mechanical Testing, Annual Book of
ASTM Standards, 2009
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(1996), 991–1005
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A. SALIMI et al.: INFLUENCE OF SAMPLE DIRECTION ON THE IMPACT TOUGHNESS ...
Materiali in tehnologije / Materials and technology 47 (2013) 3, 385–389 389

T. MAUDER et al.: EXPERIMENTAL AND NUMERICAL INVESTIGATION OF AN AIR-PCM HEAT-STORAGE UNIT
EXPERIMENTAL AND NUMERICAL INVESTIGATION
OF AN AIR-PCM HEAT-STORAGE UNIT
EKSPERIMENTALNA IN NUMERI^NA PREISKAVA ENOTE ZRAK –
PCM ZA SHRANJEVANJE TOPLOTE
Tomas Mauder, Pavel Charvat, Milan Ostry
Brno University of Technology, Faculty of Mechanical Engineering, Technicka 2896/2, Brno, Czech Republic
ymaude00@stud.fme.vutbr.cz
Prejem rokopisa – received: 2012-08-27; sprejem za objavo – accepted for publication: 2012-11-14
The phase-change materials (PCMs) are quite promising heat-storage media for applications that require a high heat-storage
capacity in a relatively narrow temperature interval. One of the areas of the latent-heat-storage application is thermal-energy
storage in solar air systems, where a lower heat-storage temperature leads to an increase in the overall efficiency of the system.
The paper deals with numerical and experimental investigations of the thermal performance of an air-PCM heat-storage unit.
The studied unit contained 100 aluminum containers filled with a paraffin-based PCM. Both experimental and numerical
investigations were done for a constant air temperature at the inlet of the heat-storage unit.
Keywords: heat exchanger, PCM, simulation
Materiali s fazno premeno (PCM) so obetavni mediji za shranjevanje toplote za uporabo, ki zahteva veliko kapaciteto
shranjevanja toplote v relativno ozkem temperaturnem intervalu. Eno od podro~ij uporabe shranjevanja latentne toplote je
shranjevanje toplotne energije v solarnih sistemih, kjer nizka temperatura shranjevanja toplote pove~uje splo{no u~inkovitost
sistema. ^lanek obravnava numeri~ne in eksperimentalne preiskave toplotne zmogljivosti enote zrak – PCM za shranjevanje
toplote. Preiskovana enota je sestavljena iz 100 aluminijevih vsebnikov, napolnjenih s parafinskim PCM. Eksperimentalne in
numeri~ne preiskave so bile izvr{ene za konstantno temperaturo zraka, pri vhodu v napravo za shranjevanje toplote.
Klju~ne besede: izmenjevalnik toplote, PCM, simulacija
1 INTRODUCTION
Most of the solar thermal systems cannot effectively
operate without thermal storage.1 Water is generally used
as a heat-storage medium in water-based solar systems,
but it is less practical for air-based systems. The rock
beds, where solid materials (usually pebbles) are used
for heat storage,2 can be used in air-based solar thermal
systems, but they have certain disadvantages. The rock
beds use a lot of space, they are difficult to clean, the
air-flow distribution in the beds is usually non-uniform
causing highly non-uniform temperature distribution in
the heat-storage medium and thus decreasing the energy
efficiency of the system. In order to achieve a higher
heat-storage density, the rock beds need to be heated up
to high temperatures leading to a decrease in the overall
energy efficiency. Promising media for thermal storage
in these applications are phase-change materials.3
The phase change of a material provides a rather high
thermal-storage capacity (and also energy-storage
density) in a narrow temperature interval around the
melting point of the material.4 Latent heat is absorbed or
released when the material changes phase from solid to
liquid or from liquid to solid. The most commonly used
PCMs are paraffins, fatty acids and esters, and various
salt hydrates.5
2 HEAT-STORAGE UNIT
PCM-based heat-storage units can be of various
designs. The studied heat-storage unit had a rather
simple design (Figure 1). It was a thermally insulated
box that contained 100 aluminum panels filled with the
PCM (arranged in 5 rows). The Rubitherm CSM panels
were used in the unit. The panels have the dimensions of
450 mm × 300 mm × 10 mm and each of them can
accommodate approximately 700 ml of PCM. The panels
were filled with the Rubitherm RT42 paraffin-based
PCM. The RT42 has a melting range from 38 °C to 43
°C, heat-storage capacity of 174 kJ/kg (in the tem-
perature range between 35 °C and 50 °C) and a thermal
conductivity of 0.2 W/(m K). The overall heat-storage
capacity of the unit was 12.3 kJ (3.4 kW h) in the
temperature interval between 25 °C and 55 °C.
In the situations when the unit is not fully charged it
can make sense to reverse the air-flow direction in the
Materiali in tehnologije / Materials and technology 47 (2013) 3, 391–394 391
UDK 536.65:519.61/.64 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(3)391(2013)
Figure 1: Schematic view of the heat-storage unit
Slika 1: Shematski prikaz enote za shranjevanje toplote
discharging mode (in comparison to the charging mode –
as indicated in Figure 1) in order to increase the outlet
air temperature and, thus, to increase the efficiency.
Another issue is the position of the PCM panels in the
unit. When the panels are positioned horizontally the
PCM in the fully melted state collects at the lower part of
the panel and there will be an air gap between the PCM
and the upper surface of the container. That gap can
significantly influence the heat transfer between the
PCM and the air passing through the heat-storage unit.
The volume change between the solid and liquid state is
rather significant for many PCMs (it can be larger than
15 %) and thus some empty space needs to be kept in the
containers to allow for that volume change. An expe-
rimental set-up was prepared in order to investigate the
thermal performance of the unit. The experimental set-up
consisted of the unit, a fan, an electric air heater, and the
temperature and air-velocity probes connected to a data
logger (Figure 2).
The experiments were carried out at a constant
air-flow rate throughout the unit and a constant inlet air
temperature.
3 NUMERICAL MODEL OF THE UNIT
The simulation tool TRNSYS 17 was used for nume-
rical investigations. TRNSYS 17 is a 1D simulation tool
that can be used for energy-performance simulations of
systems and buildings. A schematic of the numerical
model of the heat-storage unit is shown in Figure 3.
Simulations were done for 5 rows of the CSM panels
with 20 panels in each row as was the case in the expe-
rimental investigations. Since there are 19 geometrically
equal air channels between the CSM panels, only one
channel was modeled. Actually, assuming the planar
symmetry, only a half of a channel with a half of the
CSM panel thickness was modeled. The numerical
model of the heat transfer including a phase change in
the PCM unit was implemented in MATLAB and
connected to TRNSYS.6
The numerical model for the heat transfer in the PCM
created in MATLAB is based on the implementation of
the 1D heat-transfer equation that includes the source of
the latent heat of the phase change:6
∂
=
∂
+
∂ ∂t
cT k
T
x
Q( ) 
2
2 (1)
where  represents the density, c denotes the heat capa-
city, k stands for the thermal conductivity, t is the time,
T represents the temperature and x is the spatial coordi-
nate. The term Q in equation (1) can be expressed as
follows:
Q H
f
t
= 
∂
∂
s
(2)
where H denotes the latent heat and fs is the solid
fraction that represents the ratio between the solid and
liquid phases. If fs = 0, the material is in the liquid state
and, therefore, only thermo-physical properties related
to the liquid state are considered. Conversely, if 1 fs = 0,
the material is in the solid state. The theoretical analysis
of solidification is based on the equilibrium with the
assumption that a complete diffusion occurs between the
solid and liquid phases. A simple premise is to assume
that the latent heat increases linearly with the tem-
perature:
f
T T
T Ts
L
L S
=
−
−
(3)
In equation (3) TS and TL represent the solidus and
liquidus temperatures, respectively. The solution of
equation (1) strictly depends on the initial and boundary
conditions. The initial condition describes the tempe-
rature distribution for t = 0:
T x t T x( , ) ( )= =0 0 (4)
T. MAUDER et al.: EXPERIMENTAL AND NUMERICAL INVESTIGATION OF AN AIR-PCM HEAT-STORAGE UNIT
392 Materiali in tehnologije / Materials and technology 47 (2013) 3, 391–394
Figure 2: Experimental set-up
Slika 2: Eksperimentalni sestav
Figure 4: Detail of the computational domain
Slika 4: Detajl ra~unske domene
Figure 3: Simplification of the storage unit for the numerical model
Slika 3: Poenostavitev naprave za shranjevanje za numeri~ni model
The numerical model takes into account the
Neumann type of boundary conditions determined by the
heat fluxes q on the surface:
−
∂
∂
= =
T
x
x f q( , ) 0 (5)
The adiabatic boundary condition ( q = 0) is used at
the plain of symmetry in the middle of the container.
A detail of the computational domain is shown in
Figure 4. The communication between the TRNSYS
model and the MATLAB model is described below.
The computations started from a given initial
temperature profile in the PCM layer (temperatures tpcm1
to tpcmn). A constant temperature across the layer was
assumed in the simulations.
The heat flux q i at the surface of the aluminum wall
of the container was obtained from the TRNSYS model.
No thermal resistance between the PCM and the
container wall was assumed and the wall temperature ts2
was considered to be equal to the temperature tpcm1. The
heat flux obtained from the TRNSYS model was then
used as an input for the MATLAB model for the PCM
layer. The MATLAB model provided a new value of the
tpcm1 that was used for calculating the heat flux in the
next time step in the TRNSYS model. A time step of 60
s was used in the TRNSYS model while the MATLAB
model used a much shorter time step of 1 s in line with
the stability condition. The stability condition for explicit
formula was used according to7.
4 RESULTS AND DISCUSSION
Figure 5 shows the comparison of experimental and
numerical results for the heat-storage rate in the unit.
The heat-storage rate was obtained from the air-mass-
flow rate and the air temperatures at the inlet and the
outlet of the storage unit. The air temperature at the inlet
of the unit in the heat-storage period was 58 °C and the
air-flow rate was 230 m3/h. Though a constant air tempe-
rature at the inlet of the unit is of rare occurrence under
real operating conditions, it is very illustrative for the
theoretical analyses. There are certain discrepancies
between the predicted and measured heat-storage rates.
The electric air heater used in the experiments had an
output of 2 kW, but since the heater needed some time to
reach that output the measured heat-storage peak is not
as sharp as in the case of numerical simulations. The
current version of the simulation model neglects the heat
loss to the surroundings. This is one of the reasons for
the increasing discrepancies over longer periods of time.
At a certain point the air-temperature difference between
the inlet and the outlet of the unit is not due to the
storage rate but due to the thermal loss of the unit.
Another reason for discrepancies is associated with the
air flow inside the unit.
Though the heat transfer in the case of the air flow
between two parallel planes is well described in the
literature, the uncertainty remains about the air-flow
rates in particular air channels. The melting of the PCM
is yet another source of uncertainties. A proper simu-
lation of the melting and solidification requires a
complex 3D numerical model that takes into account the
convection in the liquid PCM as well as the PCM volu-
me change (to address the void formation during soli-
dification).
The heat-release rates can be seen in Figure 6. The
air temperature at the inlet of the unit was 25 °C during
the heat-discharge period. The air-flow rate was the same
as in the heat-storage period. Again, there is a relatively
good match between the numerical and experimental
results at the beginning of the heat-release period, but the
discrepancies increase with the elapsed time. The heat-
release rate peaks at more than 2 kW at the beginning of
the heat-release period, but it very quickly drops to
around 1 kW when the sensible heat above the melting
range is released. The differences between the numerical
and experimental results can be explained with the
non-uniform air-flow rates in the air channels between
the panels. The numerical model assumes that the
air-flow rates in all the air channels are the same and,
therefore, the heat fluxes stored to, or released from, all
the panels in one row are the same.
Materiali in tehnologije / Materials and technology 47 (2013) 3, 391–394 393
T. MAUDER et al.: EXPERIMENTAL AND NUMERICAL INVESTIGATION OF AN AIR-PCM HEAT-STORAGE UNIT
Figure 6: Heat-release rates
Slika 6: Hitrost spro{~anja toplote
Figure 5: Heat-storage rates
Slika 5: Hitrost shranjevanja toplote
5 CONCLUSION
A 1D simulation model of a heat-storage unit with a
PCM for thermal-energy systems using air as a heat
carrier was developed. The model of the unit was created
with the use of coupling between the TRNSYS 17
simulation tool and the in-house MATLAB model for
PCMs. An experimental setup with a test heat-storage
unit was put together in order to validate the developed
numerical model experimentally. The comparison of
experimental and numerical results revealed certain
discrepancies in the predicted and measured tempe-
ratures and energies. The simulation model neglected the
heat loss to the ambient environment, the uncertainties
associated with the air-flow inside the unit and it used a
rather simple approach to the model phase change of the
PCM. The necessary communication between TRNSYS
and MATLAB also significantly slowed down the
simulations. To address all these issues a more advanced
model of a heat-storage unit implemented as a TRNSYS
type (programmed in C++) is under development.
Acknowledgement
The authors gratefully acknowledge the financial
support from the project GACR P101/11/1047 founded
by the Czech Science Foundation, project ED0002/01/01
– NETME Centre, and the Specific Research Project
FSI-J-12-22.
6 REFERENCES
1 M. Muhieddine, É. Canot, R. March, Various approaches for solving
problems in heat conduction with phase change, International Journal
on Finite Volumes, 6 (2009), 20
2 S. Jaber, S. Ajib, Optimum design of Trombe wall system in Medi-
terranean region, Solar Energy, 85 (2011), 1891–1898
3 D. L. Zhao, Y. Li, Y. J. Dai, R. Z. Wang, Optimal study of a solar air
heating system with pebble bed energy storage, Energy Conversion
and Management, 52 (2011), 2392–2400
4 P. Pinel, C. A. Cruickshank, I. Beausoleil-Morrison, A. Wills, A
review of available methods for seasonal storage of solar thermal
energy in residential applications, Renewable and Sustainable Ener-
gy Reviews, 15 (2011), 334–3359
5 Y. Dutil, D. R. Rousse, N. Ben Salah, S. Lassue, L. Zalewski, A
review on phase-change materials: Mathematical modeling and simu-
lations, Renewable and Sustainable Energy Reviews, 15 (2011),
112–130
6 L. Klime{, P. Popela, An Implementation of Progressive Hedging
Algorithm for Engineering Problems, In: Mendel 2010, Proceedings
of the 16th International Conference on Soft Computing, Brno, BUT,
2010, 459–464
7 R. Tavakoli, P. Davami, Unconditionally stable fully explicit finite
difference solution of solidification problems, Metallurgical and
Materials Transactions B, 38 (2007) 1, 121–142
T. MAUDER et al.: EXPERIMENTAL AND NUMERICAL INVESTIGATION OF AN AIR-PCM HEAT-STORAGE UNIT
394 Materiali in tehnologije / Materials and technology 47 (2013) 3, 391–394
M. CHABI^OVSKÝ, M. RAUDENSKÝ: EXPERIMENTAL INVESTIGATION OF A HEAT-TRANSFER COEFFICIENT
EXPERIMENTAL INVESTIGATION OF A
HEAT-TRANSFER COEFFICIENT
PREISKAVE KOEFICIENTA PRENOSA TOPLOTE
Martin Chabi~ovský, Miroslav Raudenský
Heat Transfer and Fluid Flow Laboratory, Faculty of Mechanical Engineering, Brno University of Technology, Technická 2, 616 69 Brno,
Czech Republic
chabicovsky@LPTaP.fme.vutbr.cz
Prejem rokopisa – received: 2012-08-31; sprejem za objavo – accepted for publication: 2012-11-22
A special piece of apparatus was developed to study the cooling of hot steel surfaces using full-cone nozzles. This apparatus
allowed the movement of the test sheet in a vertical direction, up and down. Experiments with different water pressures and flow
rates were conducted. It was observed that the value of the average heat-transfer coefficient increased with an increase in the
water pressure for all the surface-temperature range and the dependence of the heat-transfer coefficient on the
water-impingement density is linear in the region of stable film boiling. A mathematical model based on a regression analysis
for predicting the heat-transfer coefficient in the region of stable film boiling was developed.
Keywords: cooling, impingement density, full cone nozzles, heat-transfer coefficient
Razvita je bila posebna naprava za preu~evanje ohlajanja vro~e povr{ine jekla s podro~jem s {obami. Ta naprava omogo~a
vertikalno pomikanje preizkusne plo~evine dol in gor. Izvr{eni so bili preizkusi z razli~nimi tlaki vode in hitrostmi pretokov.
Opa`eno je bilo, da nara{~a srednja vrednost koeficienta prenosa toplote z nara{~anjem tlaka vode za vsa obmo~ja temperature
povr{ine in da je linearna odvisnost med koeficientom prenosa toplote in zmanj{evanjem gostote vode v obmo~ju stabilne tanke
plasti pri vrenju. Na osnovi regresijske analize je bil razvit matemati~ni model za napovedovanje koeficienta prenosa toplote v
obmo~ju stabilne tanke plasti pri vretju.
Klju~ne besede: ohlajanje, zmanj{evanje gostote, podro~je s {obami, koeficient prenosa toplote
1 INTRODUCTION
The spray cooling of vertically moving, hot, stainless-
steel sheet was studied in the Heat Transfer and Fluid
Flow Laboratory. The temperature range from 900 °C to
200 °C was covered by these experiments. The experi-
ments with different water-impingement densities were
conducted. The dependence of the heat-transfer coeffi-
cient (HTC) on different water-impingement densities
(mL) has been studied by many authors, but it is known
that different nozzles and test conditions provide
different results. It is also known that in the range of
stable film boiling, the heat-transfer coefficient is inde-
pendent of the surface temperature and that it is mostly
influenced by the water-impingement density. It was
shown in1 (Figure 1), that the dependence of the
heat-transfer coefficient on the impingement density is
linear for water temperatures under 20 °C and
water-impingement densities of 100–2000 kg /(m2 min).
2 EXPERIMENTAL PROCESS
An experimental apparatus developed for the vertical
moving of a hot-test, stainless-steel sheet was used in the
experiments (Figure 2). The hot-test sheet with a thick-
ness of 1 mm moved vertically up and down, and it was
cooled by three rows of collectors with full-cone nozzles.
The collectors were connected through a pump to a water
tank with an adjustable water temperature. Each collec-
tor was directly connected to a manometer. The heater,
which heated the test sheet to 900 °C, was on the top of
the experimental apparatus. Four thermocouples of type
Materiali in tehnologije / Materials and technology 47 (2013) 3, 395–398 395
UDK 536.2:519.233.5 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(3)395(2013)
Figure 1: Heat-transfer coefficient as a function of the impingement
density1
Slika 1: Koeficient prenosa toplote v odvisnosti od zmanj{evanja
gostote1
K were welded on the rear side of the test sheet. The
distance between the thermocouples was 20 mm (Figure
3). The holder of the test sheet was equipped with a
position sensor and data logger, which recorded the
information about the test-sheet temperature and the
position. The test sheet was heated in a furnace at a
temperature of 900 °C and then moved up/down under
spraying nozzles until it was cooled to a temperature of
200 °C. The movement of the test sheet was conducted
with a velocity of 3 m s–1. The distance between the
nozzles was 40 mm, while the distance between the
collectors was 330 mm, and the distance from the nozzle
orifice to the test sheet was 250 mm. An example of the
nozzle configuration and photographs of the experiment
are shown on the Figure 4. The experiments were made
with a water temperature of 40 °C and differed only in
terms of the pressures. The pressures and water-impinge-
ment densities used in experiments are shown in Table 1.
Table 1: Measured experiments
Tabela 1: Rezultati meritev preizkusov
Experiment Water pressurep/bar
Water impingement
density
(kg m–2 s–1)
1 0.2 3.3
2 1.3 9.7
3 4.3 18.8
4 6 22.6
3 RESULTS
The measured temperatures were recomputed to the
surface temperatures (sprayed/cooled side) for the
location of the temperature sensors by the inverse task.2
An example of the four computed surface temperatures
is shown in the Figure 5. The heat-transfer coefficient
was then computed using the inverse conduction algo-
rithm.3 The average heat-transfer coefficient (average
value over the length of the cooling section, i.e., 660
mm) was computed for all the measured experiments.
Their comparison is shown in the Figure 6. It is evident
M. CHABI^OVSKÝ, M. RAUDENSKÝ: EXPERIMENTAL INVESTIGATION OF A HEAT-TRANSFER COEFFICIENT
396 Materiali in tehnologije / Materials and technology 47 (2013) 3, 395–398
Figure 2: Experimental apparatus
Slika 2: Preizkusna naprava
Figure 3: Rear side of the test sheet with marked positions of the
thermocouples
Slika 3: Zadnja stran preizkusne plo~evine z ozna~enimi polo`aji
termoelementov
that the average heat-transfer coefficient is increasing
with the increasing water pressure (impingement
density) and the Leidenfrost point increases with the
increasing water-impingement density. The average
heat-transfer coefficient is increasing with the decreasing
of the surface temperature in the region of the film
boiling. It is caused by thinning of the vapor layer.
Lowering the surface temperature under the Leidenfrost
temperature is connected with a change in the type of
boiling. Film boiling is changed into transition boiling
and this change in the type of boiling leads to a sharp
increase in the heat-transfer coefficient. It was also
observed that the dependence of the heat-transfer
coefficient on the water-impingement density is linear in
the region of stable film boiling (Figure 7). The first
derivative of the regression functions is increasing with
the lowering of the surface temperature (Table 2). The
constant in the regression functions represents the
measured natural convection and radiation. The result for
a surface temperature of 900 °C is in good agreement
with the result presented in1 (Figure 7). The regression
functions presented in the Table 2 can be used to
extrapolate or interpolate the average heat-transfer
coefficient for different water-impingement densities in
the region of stable film boiling. This model can only be
used for predicting the heat-transfer coefficient for para-
meters (type of nozzles, water temperature, configuration
of nozzles, etc.) that were used in the experiments from
which the regression analysis was made.
Table 2: Regression functions
Tabela 2: Regresijske funkcije
Surface temperature, T/°C Regression function
600 HTC = 212 mL +128
700 HTC = 166 mL +138
800 HTC = 123 mL +149
900 HTC = 113 mL +159
4 CONCLUSION
Experiments with horizontally oriented nozzles and
vertically moving hot sheet were conducted. Different
water-impingement densities were tested. It was
observed that the heat-transfer coefficient increases with
an increasing water-impingement density at all the
M. CHABI^OVSKÝ, M. RAUDENSKÝ: EXPERIMENTAL INVESTIGATION OF A HEAT-TRANSFER COEFFICIENT
Materiali in tehnologije / Materials and technology 47 (2013) 3, 395–398 397
Figure 7: Regression functions for predicting HTC
Slika 7: Regresijske odvisnosti za napovedovanje HTC
Figure 5: Example of computed surface temperature
Slika 5: Primer izra~unane temperature povr{ine
Figure 4: a) Nozzles configuration, b) spraying nozzles, c) nozzles
spraying the hot test sheet
Slika 4: a) Razporeditev {ob, b) razpr{ilne {obe, c) {obe, ki {kropijo
vro~o preizkusno plo~evino
Figure 6: Comparison of experiments
Slika 6: Primerjava preizkusov
surface temperatures and the Leidenfrost point increases
with an increasing water-impingement density. The
dependence of the heat-transfer coefficient on the water-
impingement density is linear in the film boiling regime.
The functions for predicting the average heat-transfer
coefficient in the surface temperature range from 600 °C
to 900 °C were determined.
Acknowledgement
The research in this paper was supported within the
project No. CZ.1.07/2.3.00/20.0188, HEATEAM –
Multidisciplinary Team for Research and Development
of Heat Processes.
5 REFERENCES
1 R. Jeschar, U. Reiners, R. Scholz, Heat transfer during water and
water-air spray cooling in the secondary cooling zone of continuous
casting plants, Proceeding of the 69th Steelmaking Conference,
Washington, 69 (1986), 551–521
2 M. Pohanka, K. A. Woodbury, A Downhill, Simplex method for
computation of interfacial heat transfer coefficients in alloy casting,
Inverse Problems in Engineering, 11 (2003), 409–424
3 M. Raudensky, Heat Transfer Coefficient Estimation by Inverse Con-
duction Algorithm, International Journal of Numerical Methods for
Heat and Fluid Flow, 3 (1993) 3, 257–266
M. CHABI^OVSKÝ, M. RAUDENSKÝ: EXPERIMENTAL INVESTIGATION OF A HEAT-TRANSFER COEFFICIENT
398 Materiali in tehnologije / Materials and technology 47 (2013) 3, 395–398
IN MEMORIAM
PROF. DR. JO@E RODI^ (26. 5. 1931–3. 3. 2013)
Od{el je prof. dr. Jo`e Rodi~. Od{el je metalurg, ki je
bil eden izmed in`enirjev, zaslu`nih za to, da je meta-
lurgija prerasla iz dobro razvite obrti v in`enirsko stroko.
Slovenska metalurgija se je razvila v pomembno teh-
ni{ko vedo v Sloveniji in postala prepoznavna v vsem
evropskem prostoru. Od{el je in`enir praktik, vodja
razvoja, ki se je hitro zavedel, da napredek tehnologije in
proizvodov ni ve~ mogo~ samo na podlagi izku{enj,
marve~ je postalo znanje temelj poznavanja metalur{kih
in tehnolo{kih procesov ter zgradbe trdnih kovin. Zapo-
slovali so ga teko~i in razvojni problemi proizvodnje in
proizvodov `elezarne na Ravnah, podpiral je tudi
raziskovanje in sofinanciranje nove raziskovalne opreme,
ki je presegalo interes podjetja. Zavedal se je razlike med
znanostjo in znanjem. Znanost nastaja v strogo nadzoro-
vanih razmerah dela laboratorijev, znanje pa je le tisti
njen del, ki ga je mogo~e vgraditi v tehnolo{ke procese,
kot jih omogo~ajo proizvodna oprema in dosegljive
surovine za izdelavo jekla. Delo dr. Jo`eta Rodi~a je {e
danes eden od temeljev proizvodnje orodnih jekel in
elektropretaljevanja pod `lindro, dveh danes vitalnih
delov podjetja Metal Ravne.
Poleg strokovnega delovanja je bil tudi odli~en peda-
gog, ki je znal izbrati {tevilne mlade sodelavce. Ti so se
razvili v vrhunske in`enirje in strokovnjake na svojem
podro~ju. Bil je izrazit zagovornik timskega dela in uva-
janja ra~unalnikov ter procesnega vodenja izdelave jekla.
Zaradi svoje komunikativnosti je bil poznan v {tevilnih
dr`avah in pridobil si je {tevilne prijatelje {iroko po
svetu. Bil je med ustanovitelji in dolgo leta tudi ~lan
uredni{kega odbora revije Materiali in tehnologije ter
zagovornik vsakoletne konference Materiali in tehnolo-
gije v Portoro`u. Kot avtor ali soavtor je objavil {tevilna
znanstvena in strokovna dela pri nas in v tujini ter na
konferencah.
Leta 1987 se je tedanji direktor Metalur{kega in{ti-
tuta v Ljubljani upokojil in dr. Rodi~ je prevzel vodenje
in{tituta. V svojem mandatu je poskrbel tudi za novo
investicijo v pilotno proizvodnjo in{tituta, to je v sistem
horizontalnega kontinuirnega ulivanja tanke `ice iz
materialov, ki se uporabljajo za navarjanje trdih plasti in
drugih specialnih zlitin. @al pa je ta tehnologija, kljub
svoji naprednosti, zaradi hitrega ni`anja cen teh mate-
Materiali in tehnologije / Materials and technology 47 (2013) 3, 399–400 399
rialov, torej iz ekonomskih razlogov, postala nekonku-
ren~na in se je nehala uporabljati.
V mladih letih je bil aktiven smu~ar, zagovornik
{porta pa je ostal vse `ivljenje. Njegova zasluga je razvoj
plavalnega {porta na Ravnah na Koro{kem. Sin in h~i sta
bila prvaka v smu~anju in plavanju ter bila ~lana dr`avne
reprezentance.
Prof. dr. Jo`e Rodi~ je bil ves ~as aktiven tudi na
podro~ju zbiranja podatkov s podro~ja lastnosti orodnih
jekel in obdelave podatkov ter njihovega sistemati~nega
urejanja. @elel si je izdelati katalog o orodnih jeklih ter
te podatkovne zbirke povezati v u~inkovit ra~unalni{ki
program za izbiro orodnih jekel. S tem se je aktivno
ukvarjal tudi {e po upokojitvi. @al pa mu je na~rte pre-
kri`ala prehitra smrt, ki ga je ustavila sredi izvajanja
svojih idej in na~rtov.
Prof. dr. Jo`eta Rodi~a se bomo spominjali kot veli-
kega strokovnjaka s podro~ja orodnih jekel, kot odprtega
in za novosti dovzetnega ~loveka ter velikega ljubitelja
svojih vnukov.
doc. dr. Matja` Torkar
Glavni in odgovorni urednik
Materiali in Tehnologije
400 Materiali in tehnologije / Materials and technology 47 (2013) 3, 399–400