VSEBINA – CONTENTS
In memoriam – Jo`ef Arh 1930–2010 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295
60 LET IN[TITUTA ZA KOVINSKE MATERIALE IN TEHNOLOGIJE LJUBLJANA /
60 YEARS OF THE INSTITUTE OF METALS AND TECHNOLOGY LJUBLJANA
60 let In{tituta za kovinske materiale in tehnologije Ljubljana
60 years of Institute of metals and technology Ljubljana
M. Jenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 297
40 let elektronske mikroanalize
40 years of electron probe mikroanalisis
F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 305
25 let vakuumske toplotne in kemotermi~ne obdelave na In{titutu za kovinske materiale in tehnologije, Ljubljana
25 years of vacuum heat treatment and surface engineering at the Institute of metals and technology, Ljubljana
V. Leskov{ek. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 307
Non-oriented electrical steel sheets
Neorientirane elektroplo~evine
D. Steiner Petrovi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 317
Effect of trace and residual elements on the hot brittleness, hot shortness and properties of 0.15–0.3 % C Al-killed steels with
a solidification microstructure
U~inek elementov v sledovih in preostalih elementov na krhkost in pokljivost v vro~em jekla z 0,15 – 0,3 % C, pomirjenega z Al in s
strjevalno mikrostrukturo
M. Torkar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 327
The influence of powder selection and process conditions on the characteristics of some oil-retaining sintered bronze bearings
Vpliv izbire prahov in pogojev izdelave na lastnosti samomazalnih sintranih le`ajev vrste Cu-Sn
B. [u{tar{i~, M. Godec, A. Kocijan, ^. Donik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 335
The influence of the microalloying elements of HSLA steel on the microstructure and mechanical properties
Vpliv mikrolegirnih elementov na mikrostrukturo in mehanske lastnosti HSLA jekla
D. A. Skobir, M. Godec, M. Balcar, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 343
Effect of 3 S reheating on charpy notch toughness and hardness of martensite and lower bainite
Vpliv 3 S pogrevanja na charpyjevo zarezno `ilavost in trdoto martenzita in spodnjega bainita
G. Kosec, F. Vodopivec, A. Smolej, J. Vojvodi~-Tuma, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 349
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Deformation mechanisms in Ti3Al-Nb alloy at elevated temperatures
Mehanizem deformacije v zlitini Ti3Al-Nb pri povi{anih temperaturah
S. Tadi}, R. Proki}-Cvetkovi}, I. Bala}, R. Heinemann-Jan~i}, K. Boji}, A. Sedmak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 357
Synthesis and characterisation of closed cells aluminium foams containing dolomite powder as foaming agent
Priprava in karakterizacija aluminijskih pen z zaprto poroznostjo, izdelanih z dolomitnim prahom kot sredstvom za penjenje
V. Kevorkijan, S. Davor [kapin, I. Paulin, B. [u{tar{i~, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 363
Imobilizacija lete~ega pepela z zasteklitvijo
Fly ash immobilization with vitrification
N. Ore{ek, F. Berk, N. Samec, F. Zupani~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 373
The impact of stagnant water on the corrosion processes in a pipeline
Vpliv zastajajo~ih voda v cevovodih na korozijske procese
M. Suban, R. Cvelbar, B. Bundara. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 379
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 44(6)293–423(2010)
MATER.
TEHNOL.
LETNIK
VOLUME 44
[TEV.
NO. 6
STR.
P. 293–423
LJUBLJANA
SLOVENIJA
NOV.–DEC.
2010
Determination of the critical pressure for a hot-water pipe with a corrosion defect
Dolo~itev kriti~nega pritiska v vro~evodni cevi s korozijsko po{kodbo
D. Kozak, @. Ivandi}, P. Kontaji}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 385
Effects of variations in alloy content and machining parameters on the strength of the intermetallic bonding between a diesel
piston and a ring carrier
Vpliv sprememb v sestavi zlitine in parametrov obdelave na trdnost intermetalne vezave med dieselskim batom in nosilcem obro~ka
A. F. Acar, F. Ozturk, M. Bayrak. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 391
Analysis of the process of crystallization of continuous cast special brass alloys with the acoustic emission
Analiza procesa kristalizacije kontinuirno lite specialne medenine z metodo akusti~ne emisije
M. Tasi}, Z. An|i}, A. Vujovi}, P. Jovani}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 397
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
The influence of adding emulsion flocculants on the effects of red-mud sedimentation
Vpliv dodatka emulzijskih flokulantov na pojave pri sedimentaciji rde~ega blata
D. Smolovi}, M. Vuk~evi}, D. Ble~i} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 403
Metallographic characterization of the joining part for EN 10216-1 and EN 10083-3 made with different consumables
Metalografska karakterizacija veznega ~lena iz EN 10216-1 in EN 10083-3, izdelanega z razli~nimi dodajnimi materiali
N. [trekelj, M. Lampi~, L. Kosec, B. Markoli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 407
LETNO KAZALO – INDEX
Letnik 44 (2010), 1–6 – Volume 44 (2010), 1–6 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 413
IN MEMORIAM – JO@EF ARH
1930–2010
V nedeljo, 12. septembra 2010, smo pospremili k
zadnjemu po~itku Jo`eta Arha, univ. dipl. ing. metalurgije
in tehnologa jeklarn v @elezarni Jesenice. Rojen je bil
16.12.1930 v Mo{njah na Gorenjskem, kjer je bil o~e
upravnik posestva Podvin. Med vojno so Nemci vas po`gali
in va{~ane izselili. Tudi dru`ina Arh je morala v izgnanstvo.
Veliko va{~anov je bilo zaposlenih v `elezarni in
verjetno je Jo`e dobil dovolj informacij, da se je odlo~il za
ta poklic. Takoj po vojni je dokon~al metalur{ko
industrijsko {olo in nato srednjo tehni{ko {olo na Jesenicah.
@elja po ve~jem znanju ga je vodila v Ljubljano na
Tehni{ko fakulteto, kjer je leta 1957 diplomiral.
Po diplomi je Jo`e Arh za~el svoj poklic v martinarni
kot asistent, kjer je izpopolnjeval svoje znanje. Delo
asistenta ga je zasvojilo do te mere, da se je v celoti posvetil
tehnologiji izdelave jekla.
V letu 1963 smo v @elezarni ustanovili raziskovalni
oddelek. Na tem oddelku so bili in`enirji, ki so delovali kot
tehnologi na podro~jih izdelave in predelave jekla, torej
in`enirji, ki so pred tem delali v proizvodnji. Z
laboratorijsko opremo smo lahko raziskovali tehnologije in
izsledke nato prenesli v proizvodnjo.
Jo`e Arh se je kot raziskovalec specializiral za tehno-
logijo izdelave vseh vrst jekel, predvsem na osnovi kroma,
niklja ter silicija za potrebe elektroindustrije. Za te namene
je imel raziskovalni oddelek visokofrekven~no pe~.
@elezarna Jesenice je programirala postavitev nove
hladne valjarne, predvsem za proizvodnjo nerjavnih jekel in
elektroplo~evin. Za investicijo je bil odobren kredit pod
pogojem, da @elezarna dobi licenco za ti dve vrsti jekla.
Izbrali smo podjetje ARMCO iz ZDA. Zato smo na
usposabljanje poslali in`enirje, med temi je bil tudi Jo`e
Arh. Usklajevanje na{e in njihove tehnologije ni bilo
vpra{ljivo, saj smo te vrste jekla `e proizvajali.
Za izdelavo mehkih in nerjavnih jekel se je Jo`e Arh ve~
let zavzemal za vakuumsko obdelavo. Vakuumiranje v
ponovci smo uvedli leta 1983 in prav tu je bil on najbolj
uspe{en in je izkoristil vse prednosti tega postopka.
Pomemben je njegov prispevek pri izbiri tehnologije v
novi elektrojeklarni, ki je za~ela obratovati leta 1987, to je
EBT-pe~ ter vakuumska naprava in naprava za kontinuirno
ulivanje.
Po reorganizaciji `elezarne je bil vi{ji raziskovalec v
ACRONI, Jesenice, do upokojitve leta 1993.
Spo{toval je strokovna mnenja drugih, s posebno skrbjo
je preverjal na{e delo z uspehi inozemskih jeklarn ter uvajal
nove postopke. Tehnika in ekonomika sta mu bili enako
pomembni, saj oboje vodi k splo{nemu napredku. Zavzemal
se je za modernizacijo, za ve~jo produktivnost ter za bolj{e
delovne razmere. Odlikovalo ga je zaupanje v lastne sile ter
doma~e znanje.
Njegovo znanje so opazili tudi v tujini. V letih od 1990
do 1998 je z nem{kim podjetjem Mannesmann Demag –
Messo Metallurgie sodeloval pri 14 podobnih projektih v
jeklarnah Evrope, Azije ter Ju`ne in Severne Amerike.
Zaradi zdravstvenih te`av se ni mogel ve~ strokovno
udejstvovati. Zadnje leto je bolezen tako napredovala, da se
mu je `ivljenje izteklo.
V @elezarni Jesenice smo imeli tednik "@elezar" in Jo`e
Arh je v njem objavljal rezultate tehnolo{kih re{itev. Njemu
so sledili tudi drugi raziskovalci do te mere, da so sklenili,
da se za~ne izdajanje "@elezarskega zbornika" v okviru
Slovenskih `elezarn. Revija se je po letu 1992 zaradi
spremenjenih pogojev izdajanja preimenovala najprej v
Kovine Zlitine Tehnologije, ~ez nekaj let pa v Materiali in
Tehnologije / Materials and Technology. Jo`e Arh je 26 let
opravljal delo glavnega urednika od ustanovitve revije do
njenega preimenovanja.
Tudi sam je napisal ve~ kot 30 strokovnih ~lankov in v
njih obravnaval vpliv elementov na lastnosti jekla,
prepihovanje taline z argonom, vpihovanje pra{kastih
materialov v talino za bolj{o ~isto~o in od`veplanje jekla,
uvajanje svin~evih jekel za obdelavo na avtomatskih
stru`nicah ter {e nekaj drugih podro~ij.
Jo`e Arh je bil aktiven v strokovni organizaciji, bil je
dolgoletni tajnik Dru{tva metalur{kih in`enirjev in tehnikov,
Jesenice. Svoje bogato znanje je prena{al tudi na dijake
Srednje tehni{ke {ole na Jesenicah. Prav ti so ga cenili kot
dobrega in uspe{nega predavatelja.
Leta 1985 je v skupini treh raziskovalcev prejel nagrado
sklada Borisa Kidri~a za izume za delo Vna{anje pra{natih
reaktantov v talino.
Za `ivljenjsko delo na podro~ju metalurgije ga je
@elezarna Jesenice nagradila s Pantzovim priznanjem za
leto 1988.
Dr. Marin Gabrov{ek,
univ. dipl. ing.
Materiali in tehnologije / Materials and technology 44 (2010) 6, 295 295

60 LET IN[TITUTA ZA KOVINSKE MATERIALE IN
TEHNOLOGIJE
60 YEARS OF THE INSTITUTE OF METALS AND TECHNOLOGY
Metalur{ki in{titut Ljubljana (MI), predhodnik
In{tituta za kovinske materiale in tehnologije je bil usta-
novljen leta 1950 z odlo~bo Sveta za kulturo in prosveto
LRS. Zaradi spremembe podro~ja delovanja se je leta
1991 preimenoval v In{titut za kovinske materiale in
tehnologije (IMT).
Temeljni kamen je bil postavljen leta 1947, ko se je
za~ela gradnja stavbe MI. Metalur{ki in{titut je bil
ustanovljen leta 1948 kot prora~unska ustanova in
sestavni del Tehni{ke visoke {ole pri Fakulteti za
rudarstvo in metalurgijo Univerze v Ljubljani. Leta 1950
je bila ustanovljena raziskovalna ustanova Metalur{ki
in{titut Ljubljana, zagon majhnega eksperimentalnega
plav`a 1. 5. 1950 pa je bil znamenje za za~etek dela.
Pobudnik ustanovitve MI in prvi direktor prof. Ciril
Rekar, nestor slovenske metalurgije, je spodbujal tesno
sodelovanje MI z oddelkom za metalurgijo Univerze v
Ljubljani. Ves univerzitetni strokovni kader je bil dopol-
nilno zaposlen na in{titutu – v povojnih letih je bilo
izrazito pomanjkanje strokovnega kadra, saj je bilo na
ozemlju celotne dr`ave samo 35 metalur{kih in`enirjev.
Vojne izgube v vrstah metalur{kih in`enirjev so bile
velike, zato bi bilo zelo nesmotrno, da bi se preostali
kader razdelil na pedago{ki in raziskovalni del. Jasno je
bilo tudi, da Univerza v Ljubljani brez raziskovalne
dejavnosti ne bo dala dobrih metalur{kih in`enirjev.
Takratni politi~ni vrh je predlagal, naj se univerza in
raziskovalna ustanova zdru`ita, tako bo imel in{titut
pomo~ v mladih, {ola pa naj bi bila povezana z meta-
lur{ko prakso in raziskovalno-razvojnim delom, da ne bi
postalo delo kabinetska u~enost; tako je ob 15-letnici
in{tituta prof. C. Rekar zapisal: »[tevilo diplomantov v
raziskovalni dejavnosti MI postaja od leta do leta manj{e
in manj pomembno. Ne le na{o metalur{ko {olo, temve~
tudi druge nove oddelke zapu{~a vsako leto desetine
metalur{kih in`enirjev. S tem pa nastaja mo`nost,
kakr{ne prej ni bilo, da more metalur{ki in{titut namestiti
in`enirje, ki imajo industrijsko prakso, in pa tiste mlade,
ki imajo nagnjenje za osnovna znanstvena raziskovanja
na metalur{kem podro~ju.«
Prav tako je prof. C. Rekar v nadaljevanju svojega
poro~ila napisal: »Sodelovanje {ole in in{tituta je korist-
no za obe strani.«
Ob praznovanju 15-letnice je imel In{titut 17 oddel-
kov s 123 zaposlenimi, od tega eno tretjino visoko{ol-
skega kadra, drugo so bili tehni~ni sodelavci in admini-
stracija; `e takrat je direktor prof. C. Rekar poudarjal, da
je premalo raziskovalnega kadra in preve~ tehni~nega
osebja, za idealno si je zami{ljal razmerje med
raziskovalci in tehni~no-administrativnim osebjem 1 : 1.
Materiali in tehnologije / Materials and technology 44 (2010) 6, 297–303 297
The Metallurgical Institute Ljubljana, predecessor of
the Institute of Metals and Technology Ljubljana, was
founded in 1950 with a decision of the Council for Culture
and Civilisation of the PR of Slovenia. However, because
of changes in the field of activity, in 1991 the name was
changed to the Institute of Metals and Technology.
The cornerstone was laid in 1947, when the con-
struction of the building of the MI was initiated. The
institute was established as a budget institution and in 1948
as part of the High Technical School at the Faculty of
Mining and Metallurgy of the University of Ljubljana. In
1950 the research was initiated with the start of work on a
small blast furnace and 1 May 1950 is the initial date of
activity.
The incentive for the founding of the MI was Prof. Ciril
Rekar, who was also the first director of Metallurgical
Institute Ljubljana and a supporter of a productive
cooperation between the MI and the Dept. of Metallurgy of
the University of Ljubljana. All the pedagogical workers
were additionally engaged part time at the institute. In the
years after the war there was a significant skills shortage, as
only 35 metallurgical engineers worked in the territory of
Slovenia. Because of the number of people lost during the
war it was thought to be unwise for the remaining number
of engineers to be distributed between the industry, peda-
gogical and research sectors. It was also clear that the
University of Ljubljana would not educate good metallur-
gical engineers in the absence of any research activities.
The political elite suggested joining the university and
research institutions. In this way, the institute would acquire
young people and the school connected the education with
metallurgical practice and development work to prevent the
change of the work to “cabinet science”. At the 15th
anniversary of the institute, Prof. C. Rekar wrote that the
number of graduates in the research activity of the institute
was becoming lower and less important as the number of
graduate engineers from the Faculty and other university
faculties increased. This created a previously nonexistent
opportunity to employ at the institute a larger number of
engineers with industrial experience and not only young
graduates with a personal inclination towards basic research
in the metallurgical field. In his report Prof. C. Rekar also
stated that the cooperation of the faculty and the institute
was useful for both sides.
At the celebration of the 15th anniversary the institute
had 123 employees, one-third of them university graduates,
and the majority of them technicians and administrative
workers. Already on that occasion Prof. C. Rekar stated
that the number of researchers was too small and that the
number of lower-school degree employees was too great.
He stated that 1:1 would be the ideal proportion between
Na podlagi temeljne uredbe je bil leta 1954 Meta-
lur{ki in{titut razgla{en za finan~no samostojen zavod s
sede`em v Ljubljani.
Z ukinitvijo fonda za napredek proizvodnje pri
Zvezni industrijski komori in med formiranjem fonda za
znanstveno delo je nastalo obdobje neurejenih dotacij in
financiranja znanstvenih zavodov, kar je spri~o stalno
nara{~ajo~ih izdatkov tudi MI nekajkrat spravilo v
neugoden finan~ni polo`aj. Udru`enje jugoslovenskih
`eljezara (UJ@) je od Izvr{nega sveta SRS prevzelo
ustanoviteljske pravice in obveznost skrbeti za MI kot
raziskovalno ustanovo `elezarske proizvodne panoge.
Leta 1973 je MI sporazumno s Sekretariatom za
kulturo in prosveto ter ustanoviteljem UJ@ podpisal
samoupravni sporazum o zdru`evanju organizacij zdru-
`enega dela v Slovenske `elezarne (S@), Ljubljana.
Ob 25-letnici MI je bil dele` financiranja od takratne
Raziskovalne skupnosti Slovenije (RSS) do 40 %, drugo
pa so prispevale projektne naloge, ki so jih naro~ile
slovenske `elezarne in kovinskopredelovalna industrija.
In{titut se je vse bolj povezoval z metalur{ko indu-
strijo, kar je bilo {e bolj izrazito v obdobju od leta 1966
do 1986, ko ga je vodil direktor Alojz Pre{ern, univ. dipl.
in`.
[ele v {estdesetih letih je bila dose`ena taka kadrov-
ska osnova, da so se lahko ustanavljali raziskovalni
oddelki v metalur{ki industriji, ki so se povezovali z
na{im predhodnikom Metalur{kim in{titutom in z
Univerzo. Delo je usmerjala in nadzirala najprej posebna
komisija, od leta 1970 dalje pa odbor za znanstveno-
raziskovalno delo, ki ga je do devetdesetih let vodil dr.
Marin Gabrov{ek. Leta 1973 se je MI vklju~il v S@. Ob
zdru`itvi je MI ohranil svoj status osrednje raziskovalne
institucije slovenske metalurgije in ga ohranil vse do
osamosvojitve Slovenije.
Leta 1969 je MI kupil prvi elektronski mikroanali-
zator, ki pomeni za~etek elektronske mikroanalize v
slovenskem in {ir{em jugoslovanskem ter vzhodno-
evropskem prostoru, kar pomeni tudi za~etek tradicije
predhodnika IMT – MI.
Neposredno povezovanje je dajalo celotni razisko-
valno-razvojni dejavnosti (RR) usmeritve, delitev dela pa
je opredeljevala vsebino in izvajalce temeljnih, aplika-
tivnih in neposrednih raziskav za re{evanje proizvodne,
kakovostne in razvojne problematike. Tako se je v
obdobju od 1986 do 1990 MI odlo~il in financiral ekspe-
rimentalno napravo za horizontalno kontinuirano litje
specialnih zlitin. Izkazalo se je, da je bila investicija v
izredno moderno in obetajo~o polindustrijsko napravo za
tisti ~as in kraj povsem neprimerna in se je in{titut komaj
izognil propadu.
Po osamosvojitvi Slovenije se je tudi slovenske
metalurgije dotaknila gospodarska kriza, saj je
proizvodnja od takratnih 800 000 ton jekla letno padla na
300 000 ton letno, zmanj{alo se je {tevilo zaposlenih na
vsega 2650. V `elezarnah so povsod iskali mo`nosti za
zni`anje stro{kov proizvodnje in kot prvo ukinili razisko-
298 Materiali in tehnologije / Materials and technology 44 (2010) 6, 297–303
the numbers of researchers and the number of technical and
administrative employees. At the 15th anniversary the
institute had 17 departments.
After a government decision on financially independent
institutions, the Metallurgical Institute became, in 1954, a
financially independent institution with its location in
Ljubljana.
With the closing of the Fund for the Progress of
Production by the Federal Industrial Chamber and during
the formation of the Federal Fund for Scientific Work a
time of unsettled conditions for financing the work of
scientific institutions was met and, because of increasing
expenses, the institute found himself in a critical financial
situation. The Union of Yugoslavian Steelworks (UJ@) took
over the funding rights and obligations from the Executive
Council of the SRS and started to finance the MI as a
research institution for the steel-producing industry.
In 1973 the MI signed, an agreement with the Secretary
for Culture and Civilisation and the UJ@ the joining to the
self-management company Slovenian steelworks Ljubljana.
At the 25th anniversary the funding of the Slovenian
Republic was at 40 %, while the difference was acquired
from projects for the Slovenian ironworks and different
manufacturing industrial companies.
The research work of the institute strengthened
constantly the connections with the metallurgical industry,
even stronger in the years 1966 to 1986, with Eng. Alojz
Pre{ern as the director.
Only in the 1960s did the number of university
graduates reach sufficient levels for the formation of
research units in metallurgical companies that established
good contacts with our predecessor and the University. The
program of work was established first by a special
commission, from 1970 by the Board for Research Work
and Prof. M. Gabrov{ek as the head. By joining with the
Slovenian steelworks, the institute conserved the status of
the central research institution of Slovenian metallurgy until
the independence of Slovenia
In 1969 an electron probe microanalyser was acquired
and the use of this advanced analytical device was started
in Slovenia, Yugoslavia and in the Eastern European area
and the germ for the quality growth of the predecessor of
today’s IMT was made possible.
Direct connections with the industrial sector gave a
general direction to the whole research activity and the
division of work determined the topics and the performers
of basic, applied and basic research projects aimed at
improving the quality of production and of products,
development and the preparation of scientific articles. In
relation to the expected development in the years 1986 to
1990, it was decided in 1990 to finance the acquisition and
construction of an experimental unit for a horizontal half
industrial continuous caster for special alloys. This decision
was later found to be a great mistake as it was not suited to
the time and place and the institute found itself in a very
critical financial situation.
After the independence of Slovenia, Slovenian
metallurgy faced a serious crisis and steel production
decreased from about 800,000 tons to about 300,000 tons

valne oddelke. Enako se je zgodilo tudi v Impolu in
Talumu, predelovalcema aluminija, propadla pa so
mnoga podjetja. S tem pa sta tako Metalur{ki in{titut kot
tudi univerza izgubila sogovornike v industriji. Nalog, ki
so jih naro~ale S@ oziroma druga metalur{ka podjetja, je
bilo vse manj in Metalur{ki in{titut si je moral najti
druge vire financiranja in se udejstvovati v drugih vejah
industrije; poleg tega pa je prenehal delovati tudi velik
jugoslovanski trg. MI je pod vodstvom takratnega
direktorja prof. Vodopivca na{el svojo ni{o predvsem v
kovinskopredelovalni industriji in v termoenergetiki.
Zaradi svoje nove dejavnosti in pritiska javnosti ter
politike, »da metalurgija nima v Sloveniji nobene prihod-
nosti«, se je Metalur{ki in{titut leta 1991 preimenoval v
In{titut za kovinske materiale in tehnologije – IMT.
Takrat se je Metalur{ki in{titut razdelil na IMT in na
proizvodni del Metalur{ki in{titut Ljubljana-pilotna
proizvodnja MIL-PP. [tevilo zaposlenih na IMT se je
zmanj{alo od 110 na 60 sodelavcev. IMT je s svojim
znanjem o materialih postal partner Nuklearne elektrarne
Kr{ko (NEK) in Termoelektrarne [o{tanj (TE[) ter
drugih slovenskih elektrarn. Raziskovalci IMT pa so se
intenzivno ukvarjali s temeljnimi in aplikativnimi
raziskavami o materialih in uspe{no konkurirali na
javnih razpisih resornega ministrstva v programu NRP
kot tudi v subvencijah Ministrstva za znanost in tehno-
logijo (MZT) gospodarskim organizacijam, kar je bil
tudi skromen za~etek ponovnega sodelovanja s slovensko
metalurgijo.
Prelomnica v zgodovini in{tituta je bilo leto 1997, ko
je postal javni raziskovalni zavod, ustanovitelja pa sta
bila MZT in z eno tretjino Slovenska industrija jeka
(SIJ). V tem obdobju so se na in{titutu ustanovili novi
laboratoriji, npr. Laboratorij za fizikalno metalurgijo
povr{in, ki je opremljen z moderno aparaturo za razis-
kavo povr{in in »in situ« prelomov kovinskih materialov
in kompozitov, Laboratorij za lezenje, v katerem so
moderne naprave, narejene na in{titutu po zgledu tujih iz
razvitega dela Evrope.
Na in{titutu je bil leta 2000 ustanovljen Laboratorij
za tlak in dve leti kasneje akreditiran kot nosilec
nacionalnih etalonov za celotno obmo~je tlaka od 10–9
mbar do 2000 bar. Leta 2000 je bila ustanovljena nova
raziskovalna skupina Vakuumistika in tehnologije.
V letu 2003 smo kupili najsodobnej{i elektronski
mikroanalizator, analitski mikroskop, prvi s Schottkyje-
vim izvirom elektronov v slovenskem prostoru in ga
naslednje leto instalirali.
V letu 2004 je in{titut ustanovil Center odli~nosti
(CO) Moderni kovinski materiali, ki je bil namenjen
prenosu znanja v industrijo in razvoju regij ter predvsem
posodobitvi raziskovalne infrastrukture. V CO je poleg
in{tituta {e pet partnerjev iz gospodarstva in javne sfere.
In{titut je nabavil visoko lo~ljiv presevni elektronski
mikroskop z energijsko disperzijskim spektrometrom za
analizo in STEM-enoto, ki omogo~a vrsti~no presevno
mikroskopijo, skupaj z napravami za pripravo vzorcev;
300 Materiali in tehnologije / Materials and technology 44 (2010) 6, 297–303
per year and the number of employees in the companies
decreased to about 2650. In the companies a reduction in
production costs was sought and one of the first steps was
the closing of the research units. At the same time the
evolution of the companies Impol and Talum, producers of
aluminium and aluminium alloys occured. A number of
other companies were closed. This was followed by a
decrease in the number of contacts of the institute and the
faculty with industrial companies and the institute was
obliged to find other financial support, as the Yugoslavian
market was also lost.
Director at that time, Prof. F. Vodopivec, found the MI
place in the metals manufacturing industry and producers
of electricity in steam power plants. Because of the change
of activity and political pressure that “metallurgy has no
future in Slovenia” in 1991 the institute changed its name
to the Institute of Metals and Technology (IMT) and
divided into two parts: IMT and the production unit
MIL-PP. The number of employees was reduced from 110
to 60. The knowledge of metallic materials enabled a
cooperation with such partners as Nuclear Power Plant
NEK Kr{ko and Steam Power Plant [o{tanj, other electric
energy-producing companies and a number of small
companies. The researchers continued their work on basic
and other topics and were successful in obtaining financing
also from the ministry responsible for research and
development. They were also associated in projects for
industrial companies in the frame of the NRP program.
This program was also the slow initial move for a new
cooperation with the Slovenian metallurgical industry.
The turning point in the institute’s history is the year
1997, when it was officially declared a public research
institution with the founders being two-thirds the Ministry
of Science and Technology and one-third the company
Slovenian Steel Industry. After this year, new laboratories
were established, for example, the Laboratory for the
Physical Metallurgy of Material Surfaces, with advanced
equipment for surfaces and the “in situ” examination of the
fracture of materials; and the Creep Laboratory with
devices of our own design and construction for the testing
of creep deformation. In 2000 the Laboratory for Pressure
was established, which 2 years later received an accredi-
tation for holding the national etalons for pressure in the
range of 10–9 mbar to 2000 bars. In 2000 a new research
group was established with the name Vacuum Science and
Technology. In 2003 a very sophisticated electron-probe
microanalyser was acquired, the fist analytical microscope
with a Schotky source of electrons in the Slovenian area,
and installed in the following year.
In 2004 the Centre of Excellence Modern Metallic
Materials was established, dedicated for the transfer of
knowledge to industry and the development of regions; first
of all, the modernisation of the research infrastructure. As
part of this centre, the institute is associated with five
partners from industry and the public sphere.
The institute acquired a high-resolution transmission
electron microscope with an energy-dispersive spectro-
meter used for analysis and an STEM device for scanning
transmission microscopy. It also acquired devices for the
kovinske materiale je namre~ potrebno stanj{ati na debe-
lino 7 nm. JEM HR 2100 je bil instaliran v letu 2008.
Tako je postal in{titut eden najmodernej{ih centrov za
raziskave strukture kovinskih materialov od povr{ine
(AES, XPS) do nano- in mikrostrukture (JSM 6500F-
FE-SEM z ED/WD in EBSD-tehnikami) ter strukture na
atomski skali (JEM 2100 HR, STEM in EDS ter z
mo`nostjo segrevanja vzorca in situ do 1000 °C) tako v
slovenskem prostoru kot v svetu.
Prav tako smo iz sredstev strukturnih skladov Evrop-
ske unije (EU) kupili najnovej{i dinami~ni 250 kN
trgalni stroj s pripadajo~o opremo za visokotemperaturne
preizkuse in s priborom za lomno mehaniko. V letu 2009
smo prenovili delavni{ko halo v Mehanski laboratorij,
kjer je instaliran nov trgalni stroj INSTRON, ter preselili
`e obstoje~ trgalni stroj 500 kN in drugo obstoje~o
opremo. Tudi na podro~ju mehanskih lastnosti kovinskih
materialov smo po opremljenosti med vodilnimi v
evropskem prostoru.
V letu 2006 je IMT pridobil akreditacijo za dva
laboratorija, in sicer za metalografski in mehanski
laboratorij ISO 17025.
In{titut ima redno, dopolnilno in pogodbeno zapo-
slenih 60 sodelavcev, od tega 35 z visoko{olsko izo-
brazbo (23 doktorjev znanosti, 1 magistra in univ. dipl.
in`.) ter vzgaja 14 podiplomskih {tudentov in 5 novih
{tudentov s Kosova ter BiH, ki so zaposleni na in{titutu
ali pa se na in{titutu usposabljajo.
In{titut se ponovno spopada s pomanjkanjem stro-
kovnega kadra, je bilo zapisano ob 50-letnici. Podobno
kot na vsej tehniki je tudi na metalurgiji in materialih
premalo {tudentov kljub kadrovskim {tipendijam in
pripravljenosti industrijskih podjetij prispevati k vzgoji
strokovnega kadra ter razpisanim {tipendijam tehnolo{ke
agencije TIA.
Zaradi pomanjkanja visoko strokovnega kadra in
zaradi nepripravljenosti Naravoslovnotehni{ke fakultete
(NTF-OMM) na enakopravno sodelovanje, se je IMT
2006 leta pridru`il Mednarodni podiplomski {oli Jo`efa
Stefana (MP[) in v {tudijskem letu 2008/2009 za~el
{tudijski program Napredni kovinski materiali (Advan-
ced Metallic Materials) v okviru smeri Nanoznanosti in
nanotehnologije. Tako je na podiplomskem izobra`e-
vanju na MP[/IMT vpisano 17 podiplomskih {tudentov,
kar je izreden uspeh in dober obet za nove vrhunske
kadre.
Julija 2010 je bil podeljen prvi doktorat, septembra
dva magisterija, v decembru pa je zagovarjal svoje dok-
torsko delo prvi podiplomski {tudent s Kosova.
In{titut si mora pridobiti do ene tretjine letnega
prora~una na trgu, torej v industriji, kar pa v dana{njem
~asu ni enostavno. Zato je na pobudo Ministrstva za
gospodarstvo za~el delovati tudi na podro~ju dr`av
zahodnega Balkana, v BiH in na Kosovu. Tako smo
realizirali projekt postavitve Laboratorija za metrologijo
tlaka, katerega odprtje je bilo junija 2009 v Sarajevu na
Meroslovnem in{titutu BiH. Projekt je financiralo Mini-
Materiali in tehnologije / Materials and technology 44 (2010) 6, 297–303 301
preparation of specimens, as it is necessary to decrease the
thickness of specimens to 7 nm. The JEM HR 2100
equipment was installed in 2008.
This development transformed the institute for the
investigation of the structures of metallic materials from the
surface (AES, XPS) to the nanodimensions of the micro-
structure (JSM 6500F- FE-SEM z ED/WD and EBSD
methods) and the structure on atomic scale (JEM 2100 HR,
STEM and EDS) and the possibility of in-situ heating of
the specimens up to 1000 °C.
With support from the EU structural funds an advanced
dynamic tensile-testing device with a capacity up to 250 kN
equipped for tests at high temperature and a determination
of the fracture toughness was acquired. In 2009 the
workshop hall was renovated and the Laboratory for
Mechanical Testing and an INSTRON 500 kN tensile
testing device were moved to the hall with the new testing
machine. Also in the area of the testing of mechanical
properties of materials IMT is positioned among the leaders
in the European area.
In 2006 IMT obtained an accreditation for the Labo-
ratories for Metallography and Mechanical Testing
according to ISO 17025.
At the institute, 60 persons are employed as regular,
complementary and contract persons. A total of 35 of them
are university graduates (23 PhDs, 1 Masters and Engs.)
and it educates 14 postgraduates and 5 new students from
Kosovo and Bosnia & Herzegovina as employees or
associates for education.
The institute meets a deficiency of professional skills,
as stated already at the 50 years jubilee. Like all areas of
technology, the institute is faced again with a scarcity of
professional skills because the number of students of
metallurgy and materials is too small, in spite of the
scholarships and industrial companies funding the
education of professional degrees in the officially published
scholarships of the Agency for Technology (TIA).
Because of the deficiency of professional skills and the
unwillingness of NTF-OMM to involve equal cooperation,
in 2006 IMT joined the Jo`ef Stefan International Post-
graduate School and started in the school year 2008/2009
with the program Advanced Metallic Materials in the frame
of the topics Nanosciences and Nanotechnology. In the
postgraduate studies Jo`ef Stefan IPS-IMT there are 17
students registered; this will result in success and great
promise for the development of excellent professional
skills. In July 2010 we had the first doctorate, in September
two Masters degrees were achieved, and in December the
first postgraduate from Kosovo obtain his doctoral degree.
The institute is obliged to find one-third of its annual
budget from external sources, and of course obtaining
funds from industry is not simple at this time. For this
reason, and with encouragement from the Ministry of the
Economy, we also started to be active in the area of the
countries of the Western Balkans, BiH and Kosovo. In the
frame of this activity the project for the foundation of the
Laboratory for Pressure Metrology opened in June 2009 in
Sarajevo at the Institute of Metrology of BiH. The project
was financed by the Ministry of the Economy from funds
strstvo za gospodarstvo RS iz sredstev dr`avne pomo~i,
za drugi projekt, prav tako za BiH, Elektri~ne veli~ine pa
smo podpisali pogodbo 31. avgusta 2010.
Na Kosovu smo soustanovili raziskovalni in{titut
R&D SHTIME, za katerega izobra`ujemo visoko stro-
kovni kader: 8 {tudentov se je vpisalo na podiplomski
{tudij MP[/IMT Advanced Metallic Materails, poleg
tega sta {e 2 {tudenta iz BiH.
Raziskovalci IMT si prizadevamo intenzivno sode-
lovati s kolegi strokovnjaki iz slovenske industrije pri
njihovih vsakodnevnih problemih, predvsem pa pri
razvoju novih izdelkov in pri optimiranju `e uvedenih
procesov in tehnologij s svojimi izku{njami in znanjem s
podro~ja temeljnih in aplikativnih raziskav. Predvsem
posku{amo animirati vodilne v podjetjih, da svoje mlade
raziskovalce usmerijo na podiplomski {tudij; na{o vlogo
pa vidimo pri usposabljanju mladih raziskovalcev (MR)
za industrijsko problematiko in delo na vrhunski
raziskovalni opremi, ki je na razpolago na in{titutu.
Metalur{ki in{titut in kasneje In{titut za kovinske
materiale in tehnologije sta ves ~as od ustanovitve dalje
uspe{no sodelovala s priznanimi mednarodnimi znan-
stvenimi institucijami; z nekaterimi je stike navezal `e
prvi direktor prof. Ciril Rekar, kot npr.: Max-Planck-
Institut für Eisenforschung, Düsseldorf, in IRSID,
Francija. Zelo dobro sodelujemo z National Institute for
Standards and Technology – NIST, ZDA, ter z vrsto
evropskih ustanov, s katerimi sodelujemo pri mednarod-
nih projektih, kot so COST, Eureka, projekt iz 4. OP
Brite Euram-Hembot in pri petih projektih iz 5. OP,
sodelujemo pri 6. OP in pri projektih Evropskega
raziskovalnega sklada za premog in jeklo – RFCS ter pri
projektih CEA – Comisariat energie Atomic.
Prav tako ne smemo pozabiti izdajateljske dejavnosti
IMT: `e leta 1992 smo prevzeli od predhodnika
ACRONI, @elezarne na Jesenicah, izdajanje @elezar-
skega zbornika. Tradicionalno revijo smo posodobili in
preimenovali v slovensko znanstveno revijo Kovine,
zlitine, tehnologije. V letu 2000 je dobila revija dana{njo
obliko in ime Materiali in tehnologije (MIT). Veliki
napori uredni{kega odbora so pripeljali do tega, da je
revija MIT citirana v SCIE od 1. januarja 2007 in v letu
2009 `e bele`ila impact faktor IF 0.143.
Izdali smo tudi tri monografije avtorjev: Leopolda
Vehovarja Korozija, Franca Vodopivca Kovine in zlitine
ter Borisa Uleta Re{ene naloge iz fizikalne metalurgije.
Poleg tega je predhodnik IMT-ja, Metalur{ki in{titut,
organiziral vsako leto Posvet o metalurgiji in kovinskih
gradivih; gonilna sila in vodja posveta je bil prof.
Vodopivec. Leta 1990 se je v organizacijski odbor vklju-
~ila M. Jenko in takrat smo za~eli prirejati konference v
hotelu Bernardin, ki je z leti prerastel v sodoben Kon-
gresni center.
@e leta 1992 sta se pridru`ila glavnemu organizatorju
konferenc IMT kot soorganizatorja Institut »Jo`ef
Stefan« in Kemijski in{titut; glavni pobudniki so bili
takratni direktor KI prof. Stane Pejovnik, prof. Drago
302 Materiali in tehnologije / Materials and technology 44 (2010) 6, 297–303
based on developmental support. The contract ratification
for the second project, for BiH Electrical quantities,
occurred on August 31 2010.
In Kosovo we cooperated in the founding of the
research Institute R&D SHTIME, for which the pro-
fessional cadres are educated: a total of 8 candidates
registered for post-graduate study at IPS/IMT Advanced
Metallic Materials, with two candidates from BiH being
enrolled in the subject.
Researchers from IMT try hard to collaborate with
colleagues from Slovenian industry by looking for
solutions to daily problems and especially in the develop-
ment of new products and improving processes and techno-
logy using our knowledge and experience from the fields of
applied and basic research. We also try to encourage the
companies to stimulate their young researchers to apply for
postgraduate study and see our role in the education of
young researchers to solve industrial problems using the
top research equipment they have at their disposal at IMT.
The Metallurgical Institute, now the Institute of Metals
and Technology, collaborates successfully with eminent
international research institutions. The initial contacts were
established by the first director, prof. Ciril Rekar, for
example, the Max-Planck-Institut für Eisenforschung
(MPIE) Germany and the Institut de Recherches de la
Siderurgie (IRSID) France. Also, there is a collaboration
with the National Institute for Standards and Technology
(NIST), USA, and some European countries with coope-
ration in international projects: COST Eureka, a project
from the 4th OP Brite Euram-Hembot, 5 projects from 5th
OP, in the 6th OP in projects of the European Fund for
Coal and Steel – RFCS and in a project with the CEA-
Commisariat a l’Energie Atomique, France.
The publishing activity of the IMT should not be
forgotten. Already in 1992 the institute took over from the
predecessor, the company ACRONI Steelwork in Jesenice,
the editing of the journal @elezarski zbornik. This tradi-
tional journal was updated and renamed as the Slovenian
scientific journal Kovine Zlitine Tehnologije – Metals
Alloys Technologies. In the year 2000, the journal acquired
its current form and the name Materiali in Tehnologije –
Materials and Technology. The editorial board succeeded
in achieving the citation of the journal in the ISI index
SCIE from January 2007 and the journal achieved the IF of
0,143.
In addition, three monographs were printed: Leopold
Vehovar Korozija (Corrosin), Franc Vodopivec Kovine in
zlitine (Metals and Alloys) and Boris Ule Re{ene naloge iz
fizikalne metalurgije (Solved problems from Physical
metallurgy).
The predecessor of IMT, the Metallurgical Institute,
organized an annual conference on metallurgy and metallic
materials. The driving force and the chairman of the
conference was Prof. F. Vodopivec. In 1990 M. Jenko
joined the organizing committee and then the conferences
were organized in the Hotel Bernardin, which with years
developed into a modern congress center.
Kolar z IJS in direktor IMT prof. Franc Vodopivec, ki so
bili tudi pobudniki za formiranje raziskovalnega polja
Materiali pri MVZT. Prof. Vodopivec je bil predsednik
programskega odbora posveta in kasneje konference vse
do svoje upokojitve 1996. leta. Posvet se je preimenoval
v Konferenco o materialih in tehnologijah in letos smo
kon~ali `e 18. po vrsti. Poleg vabljenih predavateljev je
najbolj atraktivna sekcija Mladi raziskovalci – MR, kjer
imajo mladi raziskovalci prilo`nost predstaviti v 10-mi-
nutnem govornem prispevku svoje delo v angle{kem
jeziku in se tako usposabljati za nastope na ve~jih
mednarodnih konferencah ter obenem tekmovati za
priznanje in simboli~no nagrado. Letos je bilo 35
prispevkov mladih raziskovalcev, med njimi sta bila tudi
dva mlada raziskovalca iz Karlove univerze v Pragi na
^e{kem ter univerze na Slova{kem. Vsa leta do sedaj je
vodil sekcijo MR prof. dr. Stane Pejovnik, sedanji rektor
Univerze v Ljubljani, kar si {tejemo v posebno ~ast.
Dose`eni uspehi, izku{nje in znanje iz preteklega
obdobja so neprecenljiva osnova in velika obveza za
prihodnost, vendar pa ne pomeni garancije za uspe{nost.
Samo neprestano prilagajanje novim razmeram tako v
slovenskem kot v svetovnem merilu in iskanja novega,
bolj{ega, z razvojem podprtega novega znanja omogo-
~ajo uspeh in napredek.
Lahko re~emo, da je v zadnjih 60. letih vse to IMT-ju
in predhodniku MI-ju dobro uspevalo na njuni razvojni
poti, na kateri so se tako izoblikovala zelo zna~ilna
obdobja, ki so obenem ogledalo razvoja metalurgije in
kovinskih materialov v slovenskem prostoru in v svetu.
Zahvala gre vsem dosedanjim direktorjem, ki so
zadnjih 60 let uspe{no vodili MI oziroma IMT: prof. C
Rekarju, ing. A. Pre{ernu, dr. J. Rodi~u, prof. F. Vodo-
pivcu, in prof. L. Vehovarju, kot tudi vsem stanovskim
kolegicam in kolegom, ki so pomagali ustvarjati napred-
ne zamisli. Posebna zahvala pa gre MVZT-ju, ARRS-u,
Slovenski industriji jekla, Acroni-ju, Jesenice, Impolu,
Slovenska Bistrica, Metalu Ravne, [tore-Steelu, Talumu,
Kidri~evo, Uniorju, Zre~e, Univerzi v Ljubljani in
Mednarodni podiplomski {oli Jo`ef Stefan, ki so poma-
gali Metalur{kemu in{titutu in nato IMT-ju na poti do
dana{njega uspe{nega in{tituta.
V prihajajo~ih letih si `elimo posodobiti {e preostale
laboratorije na in{titutu in jih opremiti z vrhunsko
raziskovalno opremo, prenoviti stavbo ter urediti odnose
z NTF-OMM, da bo omogo~eno sodelovanje z univerzo
pri vzgoji strokovnega kadra v duhu misli nestorja
metalurgije prof. Rekarja: »Sodelovanje {ole in in{tituta
je koristno za obe strani« ter {e uspe{nej{ega sodelovanja
z industrijskimi podjetji, predvsem v slovenskem
prostoru. Poleg tega `elimo vklju~iti mlaj{e raziskovalce
v u~ni proces na vrhunskem podiplomskem izobra`e-
vanju na Mednarodni podiplomski {oli Jo`ef Stefan,
pridobivati nove {tudente, ki bodo baza za vrhunski
strokovni kader na podro~ju kovinskih materialov.
Monika Jenko
direktorica
Materiali in tehnologije / Materials and technology 44 (2010) 6, 297–303 303
Already in 1992 the Jo`ef Stefan Institute and the
National Institute of Chemistry joined IMT, with Prof. S.
Pejovnik from NIC, Prof. D. Kolar from JSI and IMT Prof.
F. Vodopivec as the organizers. They suggested the
formation of the research field Materials (MVZT). Prof. F
Vodopivec was also chairman of the program board of the
conference up until his retirement in 1996. The conference
was renamed as the Conference on Materials and
Technology, and this year the 18th in the series was
organized. Besides the invited speakers, the most attractive
section is that for Young Researchers, where young
researchers have the opportunity to present in 10 minutes
oral presentations the results of their work in English and,
in this way, acquire experience for talking at larger
international conferences. Additionally, they compete for
recognition and a symbolic award. This year, 35 young
researchers presented their work, with two coming from the
Carl University in Prague and from the technical University
in Slovakia. So far, Prof. S. Pejovnik chaired the session for
Young Researchers, which he also did this year, in spite of
his tremendous obligations as the rector of the University of
Ljubljana. His presence and work were a special honour for
the conference.
The results obtained, the experience, and the knowledge
are a valuable base and an obligation for the future.
However, they are not a guarantee of success. Only the
steady accommodation to the new conditions in Slovenia
and the world and the search for new and better, as well as
the development of new and open knowledge, may ensure
our success and progress.
In the past 60 years IMT and its predecessor have had
success and have mirrored the development of metallurgy
in Slovenia and in the World. Thanks are due to all the past
directors of the MI and IMT: Prof. C. Rekar, Eng. A.
Pre{ern, Dr. J. Rodi~, Prof. F. Vodopivec and Prof. L.
Vehovar as well as to all the colleagues that helped us to
realize progressive ideas. Special thanks are due also to
MVZT, ARRS, the companies SIJ (Slovenian steel
industry), ACRONI Jesenice, Impol Slovenska Bistrica,
Metal Ravne, [tore-Steel, Talum, Kidri~evo, Unior Zre~e
and the University of Ljubljana that assisted the institute in
the development of today’s successful institute.
In future years we will also update the other institute
laboratories and provide them with advanced, high-quality
research equipment, to renovate the building and to
improve the relations with NTF-OMM to improve the
cooperation with university in the education of a pro-
fessional human resources in the spirit of Prof. C. Rekar:
“The cooperation of the institute and the school is useful
for both sides” and a better success in the cooperation with
industrial companies, especially in the Slovenian area. We
would like, also, to include young researchers in the
topmost postgraduate education process at Jo`ef Stefan
International Postgraduate School, to acquire new students
as a basis of the future experts for the field of metallic
materials.
Monika Jenko
director

40 LET ELEKTRONSKE MIKROANALIZE
40 YEARS OF ELECTRON PROBE MIKROANALISIS
Po letu 1965, ko se je v Sloveniji za~el hitrej{i razvoj
jeklarske tehnologije, je za~ela hitreje nara{~ati potreba
po ve~jem poznanju procesov in reakcij, od katerih sta
bili odvisni kakovost in cena metalur{kih proizvodov.
Pokazala se je potreba po napravi, ki bi omogo~ala
dovolj natan~no kemijsko analizo mikrometrskih
sestavin mikrostrukture razli~nih zlitin, npr. nekovinskih
vklju~kov v jeklih in segregacij, njihove odvisnosti od
temperature nastanka in od temperaturnega re`ima
procesa pretvorbe litih v tr`ne proizvode. Raziskovalno
okolje na tedanjem Metalur{kem in{titutu (MI) in v delu
industrije je bilo nekoliko seznanjeno z mo`nostmi nove
naprave, saj so bili leta 1958 na mednarodni konferenci v
Portoro`u, ki so jo priredili MI, Max Planck Institut für
Eisenforschung iz Düsseldorfa, Nem~ija, in Institut de
Recherches de la Siderurgie iz mesta St. Germa-
nin-en-Laye, prvi~ zunaj Francije predstavljeni rezultati
analize segregacij, ki so bile izvr{ene na taki napravi.
Naprava na tem in{titutu je bila tudi uporabljena pri
raziskovanju, ki je bilo potrebno za pripravo doktorata in
magisterija dveh raziskovalcev MI. To je bila zadostna
podlaga za odlo~itev Odbora za raziskave pri tedanjem
SOZD Slovenske `elezarne, da se za MI nabavi elek-
tronski mikroanalizator. K projektu je bilo pritegnjenih
{e nekaj in{titutov, nekaj sredstev pa je prispeval tudi
tedanji Kidri~ev sklad, in v za~etku leta 1969 je na MI
za~ela delo prva elektronska mikrosonda na Balkanu in
med prvimi v srednji Evropi.
Interes za razli~ne analize je hitro rastel med
raziskovalnimi in{titucijami in industrijskimi podjetji iz
cele Jugoslavije, in v prvih petih letih dela je bilo vsako
leto za analize vzorcev, ki niso bili iz MI, porabljenih
povpre~no 1200 ur. Analize so bile izvr{ene za
metalur{ka podjetja s podro~ij jekla, sive litine, zlitin
aluminija, bakra in dragih kovin, zlitin za elektroniko in
zobno protetiko, razli~nih stekel in kerami~nih
materialov, pa tudi za raziskovalce, ki so delovali na
podro~jih geologije, kemije in fizike trdnega stanja.
Intenziven razvoj metodologije je omogo~il, da se je
dosegla analitska lo~ljivost, ki je bila ve~ja od tiste, ki jo
je zagotavljal dobavitelj naprave, npr. 0,001 % S v trdni
raztopini v feritu, dovolj natan~na analiza vsebnosti
natrija v steklu, ~eprav se je zaradi segrevanja pod
udarom curka elektronov koncentracija tega elementa v
to~ki analize hitro zmanj{evala in to~na analiza
materialov, ki niso prenesli udara majhnega curka
elektronov.
Prej kot leto po za~etku dela naprave so bila v
znanstveni in strokovni periodiki tiskana prva dela s
podlago na rezultatih elektronske mikroanalize, kako leto
kasneje pa tiskana tudi v periodiki v drugih dr`avah.
Materiali in tehnologije / Materials and technology 44 (2010) 6, 305–306 305
After 1965, in Slovenia the development of steel
technology became faster and a better knowledge of the
processes and reactions that affected the quality of
metallurgical products became more urgent. There was a
need for a device that would make possible an accurate
chemical analysis of micrometer-size details of the micro-
structure of different alloys, for example, non-metallic
inclusions in steels, segregations in steels, aluminium and
copper alloys, the investigations of their behaviour during
the processing of cast products to final products, the under-
standing of the effects of composition, temperature and
time on the microstructure and properties and a sufficient
knowledge of the basic problems of materials science and
technology. The research and development employees in
the Metallurgical Institute (MI) and in part of industry were
given some information about the possibilities of a new
developed device, as in 1958 in Portoro` an international
conference was organised by the institutes MI, Max Planck
Institut für Eisenforschung (MPIE), Düsseldorf, Germany
and Institut de Recherches de la Siderurgie (IRSID), St.
Germanin-en Laye, France and the results of the analyses
of the segregations in steel were presented for the first time
outside of France. The device in IRSID was also used in
the investigations carried out for a doctoral and a master
thesis of two MI scientists. These were solid reasons for the
decision of the research board of the company SOZD
Slovenian Ironsworks to install at the MI an electron probe
microanalyser. The project was joined by some institutes,
and financial contributions were also obtained from the
public Kidri~ Foundation and in the beginning of 1969 the
first electron-probe microanalyser in the Balkans and one
of the first in Central Europe started to work at the MI.
The interest for analyses increased rapidly in research
institutions and industrial companies from the whole of
Yugoslavia. In the first five years of analytical work, on
average 1200 hours were used for analyses outside the
mother institute by different research institutions and
industrial companies involved in research or producing
steels, grey irons, copper and aluminium alloys, precious
alloys, alloys for electronics and dental prosthetics,
different glasses and ceramic materials and for scientists
active in the topics of geology, solid-state chemistry and
physics and dental prosthetics. The rapid development of
analytical methodology allowed us to achieve a very high
chemical sensitivity (0.001 % S in solution in ferrite), an
accurate analysis of sodium in some glasses in spite of the
local decrease of content caused by the electron analytical
beam impacting on the examined point and the proper
analysis of materials with a low stability under the impact
of the electron beam. Already after the first year of use, the
first articles with the results of the electron-probe analyses
were published in national scientific and professional
journals and some years later in foreign journals, also.
Brez pretiravanja lahko trdimo, da je ta analiza
omogo~ila tudi pospe{en razvoj proizvodne tehnologije
jekla, bakrovih in aluminijevih zlitin, rast kakovosti
proizvodov teh industrij in razvoj novih proizvodov;
pripeljala je do izvirnih dose`kov raziskovalnih in tehno-
lo{kih spoznanj, npr.: kinetika homogenizacije segregacij
v jeklih, v bakrovih in aluminijevih zlitinah, natan~na
sestava nekovinskih vklju~kov v jeklih, kvazibinarni
ravnote`ni sistemi, narava in vzroki po{kodb jeklenih
cevi v termocentralah, nov mineral, izvirna spoznanja o
bolezenskih po{kodbah zob in druga.
Utemeljen je sklep, da je za~etek uporabe elektronske
mikrosonde pospe{il tehnolo{ki razvoj na podro~ju
metalur{ke industrije, pozitivno vplival tudi na kera-
mi~ne tehnologije in omogo~il nova znanstvena spozna-
nja, ki so raziskovalce iz Slovenije uveljavila tudi v
mednarodnem prostoru. Na~in nabave in dela elektron-
ske mikrosonde bi bil lahko vzor, kako je mogo~e novo
raziskovalno napravo nabaviti, razviti metodologijo dela
in jo izkoristiti v majhnem okolju, da imata od nje korist
raziskovalna in industrijske sfera.
Franc Vodopivec
IMT Ljubljana
306 Materiali in tehnologije / Materials and technology 44 (2010) 6, 305–306
It can be concluded, without any overstatement, than
electron-probe microanalysis helped the rapid development
of the industrial technology of steel, copper and aluminium
alloys, the growth of the quality of products from these
metallic materials and also allowed us to obtain original
findings of international scientific and technological im-
portance, for example, the kinetics of the homogenisation
of segregations in steels and copper and aluminium alloys,
the accurate composition of non-metallic inclusions in
steels, the diffusion rate in the Ni-W binary system, the
quasi-binary equilibrium systems of iron with VC, TiC and
NbC carbides, the nature and the cause of failures of high-
pressure steel tubing in thermal power works, the discovery
of a new mineral, original findings on teeth lesions, etc.
The conclusion that the results of the electron-probe
microanalyser in the Metallurgical Institute after the year
1969 had a beneficial effect in Slovenia on the development
of the processes of metallic materials technology, pro-
duction, hot working and heat treatment, an increase of the
products' homogeneity and quality, also had a positive
effect on the development of ceramics and enabled us to
achieve original scientific findings that were important for
different materials and their industrial use and other
purposes is therefore justified. The results of these analyses
helped to increase the reputation of Slovenian scientists
abroad, helped also to find foreign cooperation in scientific
research and had a positive effect on publishing their results
in international scientific journals. The way used to acquire
the microanalyser could be an example of how an expen-
sive analytical device can be acquired and used in small
scientific communities.
Franc Vodopivec
IMT Ljubljana
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE
OBDELAVE NA IN[TITUTU ZA KOVINSKE MATERIALE
IN TEHNOLOGIJE, LJUBLJANA
25 YEARS OF VACUUM HEAT TREATMENT AND SURFACE
ENGINEERING AT THE INSTITUTE OF METALS AND
TECHNOLOGY, LJUBLJANA
Vojteh Leskov{ek
In{titut za kovinske materiale in tehnologije, Lepi pot 11, 1000 Ljubljana, Slovenija/
Institute of metals and technology, Lepi pot 11, 1000 Ljubljana, Slovenia
vojteh.leskovsek@imt.si
Prejem rokopisa – received: 2010-05-24; sprejem za objavo – accepted for publication: 2010-09-27
Opisan je pregled petindvajset let delovanja na podro~ju vakuumske toplotne in kemotermi~ne obdelave orodnih in hitroreznih
jekel. Raziskovalni dose`ki ob podpori drugih sodobno opremljenih laboratorijev In{titutu za kovinske materiale in tehnologije
z inovativnim na~inom omogo~ajo spremljanje in uvajanje najsodobnej{ih postopkov vakuumske toplotne obdelave in
in`enirstva povr{in ter njihovo karakterizacijo, ki je relevantna za slovensko industrijo. Pri tovrstnih tehnologijah je pomembno
znanje, {kodljiv vpliv na okolje je ni~en, poraba energije in dragih materialov je majhna, zato je raziskovalno- -razvojna
dejavnost na podro~ju vakuumske toplotne in kemotermi~ne obdelave orodnih in hitroreznih jekel v povezavi z orodjarsko,
kovinsko predelovalno, jeklarsko in avtomobilsko industrijo zelo primerna za Slovenijo, ki ima na tem podro~ju tradicijo in
usposobljene kadre.
Klju~ne besede: vakuumska toplotna obdelava, in`enirstvo povr{in, orodna jekla, hitrorezna jekla, lomna `ilavost
This review describes twenty-five years of operation in the field of vacuum heat treatment and surface engineering of tools and
high-speed steels. Research achievements, with the support of other well-equipped laboratories, enables the Institute of Metals
and Technology, an innovative approach to monitoring and for the introduction of procedures in the field of vacuum heat
treatment and surface engineering and their characterization, which is relevant to Slovenian industry. When such technology is
important knowledge, the adverse environmental impact is nil, the energy consumption and the use of expensive material is
small, from these points of view the R & D activities in the field of vacuum heat and surface engineering of tools and high-speed
steels in conjunction with the tool industry, metal processing and automotive industries is very suitable for Slovenia, which has a
tradition in this area as well as the trained staff.
Key words: vacuum heat treatment, surface engineering, tool steels, high-speed steels, fracture toughness
1 RAZVOJ TEHNOLOGIJ VAKUUMSKE
TOPLOTNE OBDELAVE IN OBDELAVE
POVR{IN S PLAZMO
Na In{titutu za kovinske materiale in tehnologije je
bil leta 1985 na pobudo prof. dr. Franca Vodopivca usta-
novljen prvi center za vakuumsko toplotno in kemo-
termi~no obdelavo (CVT&KTO) orodnih in hitroreznih
jekel v Sloveniji. Investicijo v takrat najsodobnej{o
vakuumsko pe~ IPSEN VTTC-324R (slika 1), ki je bila
podlaga za tehnolo{ke raziskave na podro~ju brez-
oglji~ne okolju prijazne in energijsko var~ne vakuumske
toplotne obdelave orodij, sta omogo~ila dr`ava in Metal
Ravne.
Karakteristike pe~i zagotavljajo optimalne mo`nosti
vakuumske toplotne obdelave najzahtevnej{ih orodij za
delo v vro~em, za delo v hladnem in orodij za brizganje
plastike.
Takrat in {e danes je ta najsodobnej{a pe~ omogo~ila,
da smo s ciljno usmerjenimi R&D projekti slovenskim
orodjarnam v najte`jem obdobju njihovega prestruktu-
riranja omogo~ili relativno hiter prehod iz konvencio-
Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315 307
1 DEVELOPMENT OF VACUUM HEAT
TREATMENT AND PLASMA SURFACE
TREATMENT TECHNOLOGY
In 1985 at the Institute of Metals and Technology, at
the suggestion of prof. dr. Franc Vodopivec, the first
centre for the vacuum heat treatment and surface
engineering (VHT & SEC) of tool steels and high-speed
steels was established in Slovenia. The investment in the
then most-advanced vacuum furnace, the IPSEN
VTTC-324R, in Figure 1, which was the basis for the
technological research in less-carbon, environmentally
friendly and energy-saving vacuum heat treatment of
tools and dies, was possible thanks to the government
and the company Metal Ravne.
The characteristics of the furnace ensure an optimal
vacuum heat treatment for the most demanding tools and
dies for hot-work and cold-work applications, as well as
for plastic injection moulding.
It has remained one of the most modern furnaces and
has offered the possibility of targeted R & D projects to
Slovenian toolmakers during the toughest period of their
UDK 621.753:539.42:621.9 ISSN 1580-2949
Pregledni ~lanek/Review article MTAEC9, 44(6)307(2010)
nalne na vakuumsko toplotno obdelavo orodij za kupce,
ki so naro~ila pogojevali z zelo visokimi kakovostnimi
standardi, ki so jim lahko zadostili le z najsodobnej{o
tehnologijo izdelave orodij.
Pomemben dejavnik, ki vpliva na trajnost orodij, je
toplotna obdelava, po kateri orodno jeklo dobi kriti~ne
lastnosti, ki pove~ajo njegovo odpornost proti prevla-
dujo~emu mehanizmu nastanka po{kodb. Optimalni
postopek toplotne obdelave je mogo~ v sodobnih
vakuumskih pe~eh s hitrim ohlajanjem v toku du{ika,
helija, vodika ali me{anice razli~nih plinov pod visokim
tlakom. Prav uvedba vakuumske toplotne obdelave je
omogo~ila, da so slovenske orodjarne ohranile svoj tr`ni
dele` in ga celo pove~ale med prilagajanjem na bolj
zahtevne nove trge, saj so zaradi relativno hitrega
prehoda na bolj{o tehnologijo toplotne obdelave, ki jo je
imel in{titut, dosegle in celo presegle zahtevano
vzdr`ljivost orodij.
Na osnovi izku{enj in dobrih rezultatov s podro~ja
vakuumske toplotne obdelave orodnih jekel smo leta
1993 na IMT-ju investirali v prvo plazemsko tehnologijo
nitriranja v pulzirajo~i plazmi, ki omogo~a kontrolirano
modificiranje povr{ine kovinskih materialov. Investicijo
je IMT-ju omogo~ilo ministrstvo s finan~no pomo~jo
podjetij KOLEKTOR, d. d., iz Idrije in ORODJARNE
GORENJE, d. o. o., iz Velenja. Reaktor za nitriranje v
pulzirajo~i plazmi METAPLAS-IONON (slika 2) je bil
izdelan po na{ih zahtevah in najnovej{ih spoznanjih o
nitriranju kompleksnih orodij in strojnih delov, izdelanih
iz jekla, titana in titanovih zlitin.
S kontroliranim modificiranjem delovnih povr{in
orodij in strojnih delov pove~amo odpornost proti obrabi,
zmanj{amo koeficient trenja, pove~amo trajno dinami~no
trdnost in korozijsko odpornost, s postopkom oksinitri-
ranja pa tudi estetski videz.
Poleg vrhunskih storitev {iroke palete kemotermi~nih
obdelav za slovenske orodjarne in kovinsko predelovalno
industrijo je reaktor primeren tudi za raziskovalno in
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
308 Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315
restructuring and allowed a relatively rapid transition
from conventional heat treatment in the vacuum heat
treatment of tools and dies for customers who required
very high quality standards that can only be met with the
latest tooling technology.
One of the most influential factors affecting the
sustainability of tools is the heat treatment of the tool
steel that ensures the critical properties, thereby
increasing the tool’s resistance to the dominant damage
mechanism for specific working operation when the tool
is subjected to stresses. The optimal heat-treatment
process is possible only in modern vacuum furnaces with
rapid cooling in a stream of nitrogen, helium or
hydrogen or in a mixture of different gasses under high
pressure. The introduction of the vacuum heat-treatment
process has enabled Slovenian toolmakers to maintain,
and actually even increase, their market share in the
critical period of adaptation to new, more demanding
markets. The produced tools reached, and even sur-
passed, the required lifetime due to the relatively rapid
transition to a better heat-treatment technology intro-
duced by the IMT.
On the basis of experience and good results in the
field of the vacuum heat treatment of tool and high-speed
steels, in 1993 IMT was the first in the country to invest
in this pulsed-plasma nitriding technology. This process
allows the controlled surface modification of metallic
materials. The investment was supported by the Govern-
ment and the companies KOLEKTOR, d. d., Idrija and
ORODJARNA GORENJE, d. o. o., Velenje. The reactor
for the pulsed-plasma nitriding METAPLAS-IONON
(Figure 2) was built according to our requirements and
the latest knowledge on the nitriding of complex tools
and dies and for steel, titanium and titanium alloy parts.
The controlled nitriding of the working surfaces of
tools, dies and machine parts increases the wear resistan-
ce, decreases the friction, increases the fatigue strength
Slika 1: Enokomorna vakuumska pe~ Ipsen
VTTC324-R s homogenim plinskim ohlaja-
njem pod visokim tlakom
Figure 1: Single-chamber Ipsen VTTC324-R
vacuum furnace with homogeneous high-pre-
ssure gas quenching
razvojno dejavnost. To je {e posebej pomembno danes,
ko so vse ve~je zahteve po trajnosti orodij za preobli-
kovanje naprednih visokotrdnostnih jekel (AHSS), za
zmanj{evanje mase strojnih delov in izbolj{anje njihovih
lastnosti. Poleg tega omogo~a tudi raziskave na podro~ju
in`enirstva povr{in, kjer je nitriranje v pulzirajo~i plazmi
osnova za t. i. dupleksni postopek (PVD ali PACVD), pri
katerem nana{amo trdo eno ali ve~plastno prevleko na
nitridno plast, ki ima dovolj veliko nosilnost.
CVT&KTO je sodeloval z ve~ kot 300 orodjarnami
in podjetji iz kovinskopredelovalne industrije iz
Slovenije in tujine. Danes delujejo v Sloveniji trije
komercialni centri za vakuumsko toplotno obdelavo, v
nekaterih tovarnah v Sloveniji pa so investirali v tovrstne
tehnologije za lastne potrebe. To je trden dokaz, da je bil
uspe{en prenos teh zahtevnih HT-tehnologij v slovenski
prostor. Zaradi dobrih izku{enj tudi v drugih podjetjih za
lastne potrebe in {e posebej zaradi okoljskih zahtev
na~rtujejo nove investicije v tovrstne tehnologije.
2 RAZISKOVALNO DELO
Z uvajanjem novih materialov, tudi nanomaterialov,
so povezani novi procesi na podlagi plazemskih tehno-
logij (npr. PACVD, PVD itd.), ki bodo v prihodnjih letih
klju~ne poleg `e uveljavljenih vakuumskih tehnologij za
kovinskopredelovalna podjetja in orodjarne, ki sodelu-
jejo z avtomobilsko industrijo in bodo hotele ohraniti
konkuren~nost. Za prenos tovrstnih tehnologij je za
vsako od novo razvitih vrst materialov potreben nov
nabor procesnih stopenj in zmogljivosti, nov na~in
karakterizacije mikrostrukture, ki omogo~a razumevanje
razli~nih lastnosti in njihove medsebojne vplive. Ta
dinami~en, interaktiven razvoj se nadaljuje in zahteva
poglobljeno razumevanje vseh dejavnikov, ki so vklju-
~eni v uvajanje novih vrst materialov in tehnologij v
uporabo. Zato je bil vsa ta leta CVT&KTO tudi razisko-
valno ploden, rezultate pa se je trudil ~im hitreje pre-
saditi v industrijsko prakso.
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315 309
and corrosion resistance and, ultimately, with oxini-
triding, the aesthetic appearance can also be improved.
In addition to a wide range of top-quality services of
thermo-chemical treatments for Slovenian toolmakers
and the metal-processing industry, the reactor is also
used for R&D. This is especially important with the
increasing focus on the sustainability of tools and dies
for the formability of advanced high-strength steels
(AHSS), reducing the weight of machine parts and
improving their properties. In addition, research activity
is possible in the field of surface engineering with the
pulsed-plasma nitriding as a basis for duplex processes
(PVD or PACVD), which make it possible to apply
single or multilayer hard coatings on a nitrided layer to
increase the bearing capacity.
By the end of 2009, VHT & SEC had serviced more
than 300 tool and metal-processing companies in
Slovenia and abroad. During this time in Slovenia three
commercial heat-treatment centres have been in opera-
tion and some companies have invested in this techno-
logy for their own needs. This expanded use confirms
that the transfer of vacuum and plasma HT technologies
in Slovenia was successful, and accordingly in other
enterprises for their own needs and environmental
requirements, in particular, new investments in these
technologies are planned.
2 RESEARCH WORK
With the introduction of new materials, nano-related
materials and new processes based on plasma techno-
logies (such as PACVD, PVD, etc.) will, in the coming
years, also become a key addition to the traditional
vacuum technologies for metal-processing companies
and tool shops, working for the automotive industry, and
wishing to remain competitive. For the transfer of such
technologies, for each of the newly developed types of
materials, a new set of procedural stages and facilities is
required and a new approach to the characterization of
Slika 2: Pe~ za nitriranje v pulzirajo~i plazmi
METAPLAS-IONON in nitrirana povr{ina dela
kolenske proteze iz Ti (desno zgoraj) in
nitriranje orodja iz jekla za delo v vro~em
(desno spodaj).
Figure 2: Furnace for pulsed-plasma nitriding
METAPLAS-IONON, the nitrided surface of a
Ti knee prostheses and the nitriding of tool steel
for a hot-work application.
Odmeven dose`ek doma in v tujini je bil razvoj
metodologije merjenja lomne `ilavosti KIc krhkih orod-
nih in hitroreznih jekel. Pri materialih z nizko duktil-
nostjo je dolo~evanje lomne `ilavosti zelo zahtevno:
predvsem je zahtevna in draga izdelava atomsko ostre
utrujenostne razpoke v korenu zareze kaljenega in
popu{~enega standardnega CT-preizku{anca. Problemi, s
katerimi se sre~ujemo pri izdelavi utrujenostne razpoke,
so odlo~ujo~e pripomogli k iskanju alternativnih preiz-
kusnih metod za dolo~evanje lomne `ilavosti. Ena od
alternativnih metod dolo~evanja lomne `ilavosti, ki smo
jo razvili na IMT, je dolo~evanje lomne `ilavosti s cilin-
dri~nimi nateznimi preizku{anci z zarezo po obodu in
utrujenostno razpoko (KIc-preizku{anec) (slika 3), ki jo
izdelamo v vrtilno-upogibnem re`imu `e pred kon~no
toplotno obdelavo 1. Ena od prednosti KIc-preizku{ancev
v primerjavi s CT-preizku{anci je, da dose`emo ravninsko-
deformacijsko stanje s preizku{anci manj{ih dimenzij 1.
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
310 Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315
microstructures for understanding the various properties
and their interrelation. This dynamic, interactive deve-
lopment continues and requires an in-depth understan-
ding of all the factors involved in introducing the use of
new types of materials and technologies. For this reason
during all those years, VHT & SEC was also active in
research and a great effort was dedicated to transplant
research results quickly into industrial practice.
A notable achievement at home and abroad was the
development of a measurement methodology for the
fracture toughness KIc of brittle tool and high-speed
steels. For materials with a low ductility the measure-
ment of fracture toughness is difficult and particularly
complex, and especially costly is obtaining the ato-
mically sharp crack in the root of the notches of standard
hardened and tempered CT specimens. The main
problem encountered in the manufacture of fatigue
cracks was the driving force in the search for alternative
test methods for the measurement of the fracture
toughness. One of such methods, developed at the IMT,
was the determination of fracture toughness with circum-
ferentially notched and fatigue-precracked tensile-test
specimens (KIc-test specimens) with the dimensions in
Figure 3. With the KIc-test specimens the fatigue crack
can be created by rotating-bending loading before the
final heat treatment. One of the advantages of such test
specimens is that plain-strain conditions can be achieved
using specimens with smaller dimensions than those of
conventional CT test specimens 1.
The advantage of the KIc-test specimens over the
standardized CT-specimens (ASTM E399-90) is the
radial symmetry, which makes them particularly suited
for studying the influence of the microstructure of
metallic materials on fracture toughness. The advantage
of these specimens is related to the heat transfer, which
provides for a completely uniform microstructure. As
already mentioned, with the KIc-test specimens the
Slika 3: Cilindri~ni natezni preizku{anec za merjenje lomne `ilavosti
z zarezo po obodu in utrujenostno razpoko v dnu zareze. Vse
dimenzije so v milimetrih.
Figure 3: Circumferentially notched and fatigue-precracked KIc-test
specimen. All dimensions are in milimeters.
Slika 4: Vpliv temperature avstenitizacije in popu{~anja na trdoto HRc
in lomno `ilavost KIc vakuumsko toplotno obdelanega jekla za delo v
vro~em H11
Figure 4: Effect of austenitizing and tempering temperatures on the
HRc hardness and fracture toughness KIc of vacuum heat-treated
hot-work steel H11
– KIc-preizku{anec: cilindri~ni natezni preizku{anec z zarezo po obodu in
utrujenostno razpoko v dnu zareze  10mm x 120 mm
– temperatura avstenitizacije: 1000 °C, 1020 °C in 1050 °C
– ~as zadr`evanja na temperaturi avstenitizacije: 20 min
– ohlajanje: v toku N2 pri tlaku 1,05 bar do temperature 100 °C
– parametri ohlajanja 800-500 : 1,04; 1,02; 1,11
– prvo popu{~anje: 2 h pri 540 °C
– drugo popu{~anje: 2 h med 540 °C in 620 °C
– KIc-test specimens: circumferentially notched and fatigue-precracked tensile
specimens  10 mm × 120 mm
– austenitization temperature: 1000 °C, 1020 °C and 1050 °C
– soaking time: 20 min
– quenching: gas quenching in N2 at a pressure of 1.05 bar to 100 °C
– cooling parameters 800-500 : 1.04; 1.02; 1.11
– first tempering: 2 h at 540 °C
– second tempering: 2 h between 540 °C and 620 °C
Prednost tak{nih preizku{ancev v primerjavi s
standardnimi CT-preizku{anci (ASTM E399-90) je tudi
njihova radialna simetrija. Zato so {e posebej primerni za
opredelitev vpliva mikrostrukture kovinskih materialov
na lomno `ilavost. Izoblikovanje mikrostrukture po
obodu je popolnoma enakomerno zaradi radialno
simetri~nega odvajanja toplote. Pri merjenju lomne
`ilavosti trdih in krhkih kovinskih materialov, kjer nam
zaradi velike zarezne ob~utljivosti le s te`avo, ~e sploh,
uspe izdelati razpoko z utrujanjem, lahko utrujenostno
razpoko na tovrstnemu preizku{ancu izdelamo {e pred
kon~no toplotno obdelavo.
To nam je omogo~ilo, da smo pri vakuumski toplotni
obdelavi orodnih in hitroreznih jekel lomno `ilavost KIc
uvedli kot drugi parameter. Tako lahko iz diagrama po-
pu{~anja za jeklo, izbrano za specifi~no operacijo,
izberemo najprimernej{e razmerje med `ilavostjo in
trdoto. Taki diagrami popu{~anja nam omogo~ajo, da
izberemo parametre vakuumske toplotne obdelave, s
katerimi to tudi dose`emo pri izbranem jeklu (slika 4).
Na osnovi poznane lomne `ilavosti lahko izra~unamo
kriti~no velikost napake, ki jo orodno jeklo pri lomni
napetosti {e prenese (slika 5).
Tudi pri raziskovalnem delu so stro{ki eksperimentov
pomembni, zato po meritvi lomne `ilavosti dele KIc-pre-
izku{ancev uporabljamo {e za meritve trdote, upogibne
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315 311
Slika 5: Vpliv temperature avstenitizacije na lomne napetosti f za
jeklo H11
Figure 5: Influence of austenitizing temperature on the fracture stress
f of H11
fatigue crack can be created with rotating-bending
loading before the final heat treatment.
This allowed us to introduce the fracture toughness
KIc as a second parameter in the vacuum heat treatment
of tool and high-speed steels. In this way, from the
tempering diagram for the selected steel it is possible to
choose the most appropriate relationship between the
toughness and the hardness for a specific application.
Such a tempering diagram allows us to also choose the
parameters of vacuum heat treatment that give the
selected steel the desired performance (Figure 4).
Knowing the fracture toughness KIc and the ultimate
tensile stress u of the investigated tool steel, the fracture
stress f and the critical defects’ size in the tool steel can
be calculated (Figure 5).
The costs of experiments for research projects are
important. For this reason, after a determination of the
fracture toughness, two parts of the KIc-test specimens
are used to manufacture different samples for the hard-
ness measurements, the four-point bending, the com-
pressive strength and strain hardening exponent determi-
nation, the analysis of fracture surfaces and of the
microstructure and for the specimens of other tribolo-
gical and technological tests, necessary for under-
standing the behaviour of tool steels and high-speed
steels in practice (Figure 6).
Slika 6: Mo`nosti izdelave razli~nih preizku{ancev iz KIc-preizku{anca
Figure 6: Possibilities for the manufacturing of various trial specimens
from the KIc-test specimens
trdnosti, tla~ne trdnosti, dolo~itev eksponenta deforma-
cijskega utrjevanja, analizo prelomnih povr{in, mikro-
strukturno analizo ter za izdelavo preizku{ancev za druge
tribolo{ke in tehnolo{ke preizkuse, ki so pomembni za
razumevanje vedenja orodnih in hitroreznih jekel v
praksi (slika 6).
Ker so vse meritve izvedene na istem preizku{ancu,
lahko za razli~ne lastnosti ugotovimo, v kak{ni medse-
bojni povezavi so posamezne lastnosti.
Odmeven znanstveni dose`ek je razvoj polempiri~ne
ena~be (1), ki omogo~a, da na osnovi trdote Rockwell-C,
volumskega dele`a zaostalega avstenita faust, volumskega
dele`a neraztopljenih evtekti~nih karbidov f
c
, srednje
razdalje med temi karbidi dp, kumulativnega dele`a
fc a≥ crit karbidov in/ali karbidnih skupkov, ki so  acrit
(acrit, kriti~na velikost napake) in modula elasti~nosti E
izra~unamo lomno `ilavost za hitrorezna in ledeburitna
jekla s trdoto HRc ve~jo od 57 HRc 2.
( )
K
Ed f fc a
Ic
p c
HRc
HRc 53
crit
= ⋅
−
⎛
⎝
⎜ ⎞
⎠
⎟ ⋅
⋅ ⋅ ⋅ −− ≥
1363
1
1
6
,
( ) ([ ][ ]) (⋅ +1 f maust MPa
(1)
Na osnovi te ena~be je mogo~e teoreti~no pojasniti
vpliv mikrostrukture hitroreznih in ledeburitni jekel na
lomno `ilavost. Namre~, s toplotno obdelavo jeklu s
spreminjanjem mikrostrukture spreminjamo lastnosti.
Za orodni jekli za delo v vro~em H11 in H13 je bila
razvita empiri~na ena~ba (2) s katero na osnovi udarne
`ilavosti Charpy-V in trdote Rockvell-C ocenimo lomno
`ilavost orodnega jekla:
[ ]K CVN mIc HRc MPa= ⋅ ⋅ −4 53 1 11 0 135, , , (2)
kjer je CVN absorbirano udarno delo in HRc trdota
Rockwell-C 3.
Na sliki 7 je prikazano ujemanje med eksperimen-
talno dolo~eno lomno `ilavostjo in lomno `ilavostjo,
izra~unano na osnovi ena~be (2) za jekli H11 in H13.
Tovrstne ocene so {e posebej pomembne pri analizi
vzrokov nastanka po{kodb orodja (npr. prelom orodja
med obratovanjem itd.), pri katerem ni mogo~a izdelava
312 Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
Since all the measurements are performed on the
same test specimen, it is possible to find the correlation
between different characteristics for individual proper-
ties.
A scientifically notable achievement was the deve-
lopment of the semi-empirical equation (1) that makes it
possible to calculate the fracture toughness of high-speed
and ledeburitic steels 2 on the basis of the modulus of
elasticity E, the mean distance between the undissolved
eutectic carbides, dp, the Rockwell-C hardness, fc and faust
as the volume fractions of undissolved eutectic carbides
and retained austenite, and fc a≥ crit as the cumulative
fraction of undissolved eutectic carbides and/or carbide
clusters equal to, or larger than, the critical defect size
( acrit), with a Rockwell-C hardness higher than 57
HRc.
( )
K
Ed f fc a
Ic
p c
HRc
HRc 53
crit
= ⋅
−
⎛
⎝
⎜ ⎞
⎠
⎟ ⋅
⋅ ⋅ ⋅ −− ≥
1363
1
1
6
,
( ) ([ ][ ]) (⋅ +1 f maust MPa
(1)
On the basis of this equation it is theoretically
possible to determine the influence of individual con-
stituents of the microstructure of high-speed and
ledeburitic steels on the fracture toughness. Namely, the
heat treatment can modify the material properties by
changing the microstructure.
For the hot-work tool steels H11 and H13, an empi-
rical equation based on the Charpy-V impact toughness
and the Rockvell-C hardness was developed (2), to esti-
mate the fracture toughness of tool steels, respectively:
[ ]K CVN mIc HRc MPa= ⋅ ⋅ −4 53 1 11 0 135, , , (2)
where CVN is the energy absorption and HRc is the
Rockwell-C hardness 3.
Figure 7 shows the correlation between the fracture
toughness of the experimentally determined and cal-
culated fracture toughness on the basis of equation (2)
for H11 and H13 steels.
Such assessments are particularly important for the
analysis of the failures for tools (e.g., tool cracking
Slika 7: Primerjava med eksperimentalno dolo~enimi vrednostmi
lomne `ilavosti KIc in izra~unanimi pri enaki trdoti za jekli H11 in H13
iz CVN udarne `ilavosti in trdote HRc z ena~bo (2)
Figure 7: Comparison between the experimentally obtained values of
KIc and the pertained hardness, and the values calculated for the
investigated hot-work tool steels H11 and H13 from CVN test results
and the pertained hardness (eq. 2)
standardnih ali nestandardnih preizku{ancev za dolo~e-
vanje lomne `ilavosti. Mo`no pa je izrezati metalografski
obrus pri hitroreznih in ledeburitnih jeklih ali CVN-pre-
izku{anec pri orodnem jeklu za delo v vro~em.
Na podro~ju modificiranja delovnih povr{in orodij z
nitriranje je pomemben razvoj metodologije merjenja
lomne `ilavosti nitridne plasti z metodo vtiskovanja
Vickersove piramide pri razli~nih obte`bah (slika 8).
Za relativno tanke krhke plasti na relativno `ilavi
podlagi, kot so razli~ne nitridne plasti na orodnem jeklu,
smo ugotovili, da je za oceno lomne `ilavosti primerna
ena~ba, ki jo je razvil Shetty 4 na osnovi pojava Palmqvi-
stovih razpok 5:
[ ]K P
al
mIc MPa=
⎛
⎝
⎜ ⎞
⎠
⎟0 0319 1 2, / (3)
kjer je P obte`ba, a je polovica srednje dol`ine diago-
nale in l l ii= =∑
1
4 0
4
je srednja dol`ina razpok 6.
Pomemben rezultat raziskav na podro~ju nitridnih
plasti je ugotovitev, da lahko na osnovi izmerjenega
profila mikrotrdote dolo~imo globino maksimalnih
tla~nih napetosti, kar omogo~a optimizacijo postopka
nitriranja, tako da globino maksimalnih napetosti dose-
`emo v podro~ju maksimalnega Hertzovega tlaka, kar
pove~uje trajno nihajno trdnost (slika 9) 7.
Podeljeni so bili tudi trije patenti, in sicer za "Induk-
cijsko ogrevano celico za toplotno in kemotermi~no
obdelavo kovin v zvrtin~eni plasti: patent SI 9800119";
"Dvostranski lamelni skobelni no` za obdelavo kovin,
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315 313
during operation, etc.) in the case when it is not possible
to manufacture standard CT or non-standard cylindrical
specimens for the fracture toughness measurement. But
it is possible to cut out the metallographic sample in the
case of high-speed or ledeburitic steels and CVN
specimens in the case of hot-work tool steel.
For the modified working surfaces of tools and dies
with nitriding, of essential importance is the method of
measuring the fracture toughness of the nitride layers
with Vickers indentations at various loads (Figure 8).
For a relatively thin and brittle layer on a relatively
tough substrate, as, for example, the different nitride
layers (Figure 8) on tool steel, we found that for the
assessment of the fracture toughness the Shetty 4
equation (3), developed on the basis of Palmqvist cracks,
is appropriate 5:
[ ]K P
al
mIc MPa=
⎛
⎝
⎜ ⎞
⎠
⎟0 0319 1 2, / (3)
where P is the load, a medium-length, half of the dia-
gonal and l l ii= =∑
1
4 0
4
the mean length of the cracks 6.
For nitride layers, it is important that the depth of the
maximum compressive stresses can be determined on the
basis of the measured microhardness profile. This makes
it possible to optimise the nitriding process and ensures
that the depth of the maximum compressive stresses is
achieved in the area of the maximum Hertzian pressure
and the obtaining of an increase in the fatigue strength
(Figure 9) 7.
Slika 8: Vickersov vtis na krhki nitridni plasti; Palmqvist-ove razpoke,
difuzijska plast, spojinska plast ’ in/ali  5
Figure 8: Vickers indentation on brittle nitride layer; Palmqvist cracks,
diffusion layer, compound layer of ’ and/or  5
Slika 9: Primerjava profila mikrotrdote, njegovega odvoda po globini
in izmerjenega profila zaostalih napetosti na nitriranem jeklu za delo v
vro~em H11
Figure 9: Comparison of microhardness depth profile, its derivation
over depth and the residual stress profile for the plasma-nitrided
hot-work tool steel H11
lesa itd., sestavljen iz nosilne lamele iz konstrukcijskega
jekla in dveh enako dolgih, vendar o`jih in tanj{ih lamel
iz kabidne trdine WC-Co: patent SI 9600034 A" in
"Pove~anje trdote povr{ine zlitine FeAl v masnih dele`ih
(%) s postopkom nitriranja v pulzirajo~i plazmi: patent
SI 960001"4.
Najnovej{e raziskave so usmerjene na podro~je
podhlajevanja orodnih jekel v teko~em du{iku in nana-
{anja trdih prevlek po dupleksnem postopku PACVD, ki
je klju~en za trajnost velikih kompleksnih orodij za
oblikovanje naprednih, visokotrdnostnih jekel (AHSS),
katerih trdnost presega 800 MPa (slika 10) 8.
Pri preoblikovanju AHSS-jekel je pomembno, da pri
orodju optimiramo sistem: trda prevleka /modificirana
povr{ina/ orodno jeklo z ozirom na kriti~ne lastnosti, tj.
trdota, razteznost, `ilavost in omogo~imo doseganje
lastnosti, ki pove~ujejo odpornost sistema proti
prevladujo~emu mehanizmu nastanka po{kodb (obraba,
plasti~na deformacija, kru{enje, prelom in lepljenje)
(slika 11).
3 SKLEP
Slovenske orodjarne in podjetja s podro~ja kovin-
skopredelovalne industrije, ki sodelujejo z avtomobilsko
industrijo, imajo relativno {iroke programe, vendar pa je
V. LESKOV[EK: 25 LET VAKUUMSKE TOPLOTNE IN KEMOTERMI^NE OBDELAVE ...
314 Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315
The following patents were awarded: for the
"Induction heated cell for heat treatment and thermo
chemical treatment of metals in fluidized bed: Patent SI
9800119", for the "Lamellar bilateral planning knife for
treatment metal, wood, etc. consisting of carrier strip of
construction steel and two equally long but narrow and
thin strips of WC-Co hard metal: Patent SI 9600034 A"
and for the "Increase the surface hardness of FeAl alloy
with the pulsed plasma nitriding: Patent SI 9600014".
Recent research has focused on the area of the
deep-cryogenic treatment of tool steels and the
application of hard coatings deposited by the PACVD
duplex process, which is crucial for the sustainability of
large and complex tools for the forming of advanced
high-strength steels (AHSS), whose strength exceeds 800
MPa (Figure 10) 8.
For the forming of AHSS steels, it is important that
the system of working surfaces, i.e., hard coating /
modified surface / tool steel of the tool are optimised
with respect to the critical properties, i.e., hardness / duc-
tility / toughness, that increase the resistance to the
dominant damage mechanism (adhesion, galling, abra-
sion, plastic deformation, chipping and cracking) (Figu-
re 11).
Slika 11: Profil mikrotrdote HV0,1 in HV0,003 ter lomne `ilavosti KIc,
nitriranega vzorca iz jekla H11 z nanosom ve~plastne trde prevleke iz
TiB2/Ti-B-N/ TiN, nanesene po dupleksnem PACVD- postopku
Figure 11: Micrograph showing HV microhardness depth profile of
duplex-treated H11 substrate at a load of 1N and 0.3 N and the
corresponding fracture toughness KIc of the TiB2/Ti-B-N/ TiN
multilayer hard coating.
Slika 10: PACVD-dupleksna ve~plastna trda prevleka TiB2/Ti-B-N (21plasti)/TiN na jeklu za delo v vro~em H11
Figure 10: Duplex multilayer PACVD hard coating TiB2/Ti-B-N (21 layers) / TiN on hot-work tool steel, H11
vse manj podro~ij, kjer lahko ohranijo ali na novo pri-
dobijo konkuren~ne prednosti. Zato se vse bolj pojavlja
potreba po prestrukturiranju in povezavah, kajti le dovolj
veliki in uspe{ni subjekti, ki so povezani s centri znanja,
lahko obvladujejo trg in zmorejo velika vlaganja v
znanje in razvoj proizvodov z ve~jo dodano vrednostjo
ter v posodobitev opreme, da obdr`ijo konkuren~nost in
zagotovijo akumulacijo, s katero pokrijejo stro{ke in
zagotovijo lastnikom zanimiv profit.
Tehnologije in metodologije, ki jih uvajamo in
razvijamo ter so podprte z ekspertnim znanjem drugih
laboratorijev na IMT, spadajo v HT-tehnologije in so
primerljive s tistimi, ki imajo od tri- do petkrat ve~jo
dodano vrednost.
Ta kratka predstavitev dejavnosti in raziskovalnih
dose`kov je dokaz, da CVT&KTO ob podpori drugih
sodobno opremljenih laboratorijev za karakterizacijo na
IMT z inovativnim na~inom omogo~a spremljanje in
uvajanje najsodobnej{ih postopkov s podro~ja
vakuumske toplotne obdelave in in`enirstva povr{in in
njihovo karakterizacijo. ^e k temu dodamo {e dejstvo, da
se pri tovrstnih tehnologijah prodaja predvsem znanje, da
je ni~en {kodljiv vpliv na okolje, poraba energije in
dragih materialov pa majhna, je logi~en sklep, da je
R&D-dejavnost v povezavi z orodjarsko, kovinskopre-
delovalno in avtomobilsko industrijo, zelo primerna za
Slovenijo, ki ima na tem podro~ju tradicijo in usposob-
ljene kadre.
4 LITERATURA/REFERENCES
1 V. Leskovsek, B. Ule, B. Liscic Relations between fracture tough-
ness, hardness and microstructure of vacuum heat-treated high-speed
steel, Journal of Materials Processing Technology, 127 (2002) 3,
298–308
2 V. Leskovsek, Modelling of High-Speed Steels Fracture Toughness,
Materials and Manufacturing Processes, 24 (2009) 6, 603–609
3 V. Leskovsek, Correlation between the K-ic, the HRc and the charpy
V-notch test results for H11/H13 hot-work tool steels at room
temperature, Steel Research International, 79 (2008) 4, 306–313
4 D. K. Shetty, I. G. Wright, P. N. Mincer, A. H. Clauer, J. Mater. Sci.,
20 (1985), 1873–1882
5 S. Palmqvist, Arch. Eisenhuttenwes., 33 (1962) 9, 629–634
6 D. Nolan, V. Leskovsek, M. Jenko, Estimation of fracture toughness
of nitride compound layers on tool steel by application of the Vickers
indentation method, Surface and Coatings Technology, 1 (2006) 1–2,
182–188
7 V. Leskovsek, B. Podgornik, D. Nolan, Modelling of residual stress
profiles in plasma nitrided tool steel, Materials Characterization, 59,
(2008) 4, 454–461
8 V. Leskovsek, B. Podgornik, M. Jenko, A PACVD duplex coating for
hot-forging applications, Wear, 266 (2009) 3–4, 453–460
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Materiali in tehnologije / Materials and technology 44 (2010) 6, 307–315 315
3 CONCLUSION
Slovenian toolmakers and companies in the metal-
processing industry working for the automotive industry
have relatively large programs, but in fewer areas can
they maintain or regain a competitive advantage.
Therefore, there is a growing need for restructuring and
connections, because only large enough and successful
entities associated with centres of knowledge can be
successful in the market and are capable of a significant
investment in knowledge and the development of
products with higher added value and updated equipment
to maintain competitiveness and to ensure the accumu-
lation, which cover the costs and provide an attractive
profit for owners.
The introduced and developed technologies and me-
thodologies are also supported by the expert knowledge
of other laboratories at the IMT, are part of HT tech-
nology and comparable to those of three to five times
higher added value.
This brief presentation of the activity and achieve-
ments shows that VHT & SEC, with the support of other
well-equipped IMT laboratories for characterization, and
an innovative approach make it possible to monitor and
implement cutting-edge procedures of vacuum heat
treatment and surface engineering and their characte-
rization. If the fact is added that in these technologies, it
is mainly knowledge that is being sold, the environ-
mental impact is practically zero and the consumption of
energy and precious material is small, the conclusion is
logical that the R & D activity related to the tool-making
industry, the metal-processing and automotive industries,
is very appropriate and promising for Slovenia, which
has trained staff and a tradition in this area.

D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
NON-ORIENTED ELECTRICAL STEEL SHEETS
NEORIENTIRANE ELEKTROPLO^EVINE
Darja Steiner Petrovi~
Institute of Metals and Technology, Lepi pot 11, Ljubljana, Slovenia
darja.steine@imt.si
Prejem rokopisa – received: 2009-12-11; sprejem za objavo – accepted for publication: 2010-10-20
Non-oriented electrical steel sheets are sheets tailored to produce specific properties and are produced from Fe-Si or Fe-Si-Al
alloys. Non-oriented electrical steel sheets are incorporated into a wide range of equipment, from the simplest domestic
appliances to hybrid and pure electric vehicles. Future efforts will be focused on controlling the residual elements in the steel,
optimizing the hot and cold rolling, and optimizing the crystallographic texture development, with the aim to enhance the
performance of the finished product.
Key words: Non-oriented electrical steel sheets, silicon steel
Neorientirano elektroplo~evino izdelujemo iz silicijevega jekla in iz enakega jekla, ki je legirano z aluminijem. Neorientirano
elektroplo~evino uporabljamo predvsem za magnetna jedra elektri~nih motorjev in transformatorjev, v zadnjem ~asu pa tudi za
izdelavo hibridnih in elektroavtomobilov. Lastnosti neorientirane elektroplo~evine lahko v prihodnje {e izbolj{amo z
zmanj{anjem vsebnosti ne~isto~ v jeklu, z optimizacijo vro~ega in hladnega valjanja ter z razvojem ugodne kristalografske
teksture. Dobre lastnosti elektroplo~evin pove~ujejo energijsko u~inkovitost elektri~nih naprav, v katere so vgrajene.
Klju~ne besede: neorientirane elektroplo~evine, silicijevo jeklo
1 INTRODUCTION
Soft magnetic materials are ubiquitous in the current
electronics-based economy. Silicon steel was developed
at the beginning of the 20th century and soon became the
preferred core material for large transformers, motors,
and generators. Silicon-bearing steels are used as soft
magnetic materials in electrical appliances and devices
and are rated in terms of power loss when magnetized in
an alternating electric field. The total amount of these
steels is around 1 % of the world production of steel1.
The Worldsteel Committee on Economic Studies,
Brussels, reports that the worldwide production of strip
in 2008 was 10,291,000 metric tons, and in the EU it was
1,498,000 metric tons. The production of electrical steel
sheet and strip in the last 10 years has almost doubled.
The production of non-oriented electrical steel in
Slovenia is approximately 100,000 metric tons per year.2
Texture is one of the most important parameters
determining the magnetic properties of steel sheets. The
ideal texture of non-oriented silicon steel sheets would
be a cubic texture with grains with their (001) or (110)
planes parallel to the plane of the sheet and a uniform
distribution of the [100] direction, whereas the Goss
texture with a (110)[100] crystallographic orientation of
the grains is the typical grain structure of grain-oriented
silicon steel.
Silicon steels are fundamental for the economy of
electrical appliances, and offer the best combination for
transmitting and distributing electrical energy. The
properties required of these steels are a high permeability
and induction, low magnetic losses, and low magneto-
striction. A high permeability and induction reduce the
size and weight of the parts; low magnetic losses
diminish the the generation of Joule heat and energy
consumption; and a low magnetostriction reduces the
noise (which appears as humming) in transformers and
high-capacity machines3.
The basic technology of production for non-oriented,
fully processed electrical steels has not changed signifi-
cantly in recent decades: the basic chemistry is similar in
terms of the main alloying elements and the processing
steps are basically unchanged. Nevertheless, the losses in
a steel with a given Si and Al content are today much
lower compared to previous decades. Accordingly,
electrical steel producers have made only very small
changes to the basic chemistry used for most commercial
standard grades. International and national standards
only specify the maximum loss (and often also a
minimization of polarization/permeability), but have in
principle no lower limit to the losses. Consequently, for a
given standardized grade the difference between the
guaranteed maximum loss and the actual loss of the
material produced has increased significantly over time.
An electrical steel is a commodity product with a market
price very much determined by its grade designation.
From the steel user’s point of view, this development has
brought advantages, but it has also increased the
variability in the market and it is uncertain as to what a
standardized grade really is.
2 CLASSIFICATION
Non-oriented electrical steel sheets, commercially
also called lamination steel, silicon electrical steel,
silicon steel or transformer steel, are special steel sheets
tailored to produce certain magnetic properties. They are
used in the form of lamination stacks, mainly in electric
Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325 317
UDK 669.14.018.583 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 44(6)317(2010)
motors, transformers and alternators, depending on their
properties.
Non-oriented electrical steel sheets can be divided
into two categories:
– fully-processed grades, which are delivered in the
finished condition, continuously annealed and
sometimes varnished. They have guaranteed magne-
tic properties, in accordance with standards,4 e.g.,
EN 10106:2009.
– semi-processed grades, that are given the final
annealing treatment to develop their magnetic
properties by the user.
Non-oriented electrical steel sheets are usually
manufactured in the form of cold-rolled sheets/strips
with thicknesses of (0.35, 0.50, 0.65 and 1.00) mm and
are classified according to the value of the maximum
specific total loss in W/kg. The non-oriented electrical
steel is supplied in stacks in the case of sheets and in
coils in the case of strips (Figure 1).
The main types of non-oriented electrical steels pro-
duced in Slovenia, by Acroni d.o.o., Jesenice, are:
– Cold-rolled, fully-processed electrical steels –
DINAMO,
– Cold-rolled, semi-processed electrical steels –
ELMAG,
– Cold-rolled, fully-processed, high-permeability elec-
trical steels – PERMAG FP
– Semi-processed, high-permeability, electrical steels
– PERMAG SP.5
The precise technologies and metallurgical processes,
combined with technical development and investments in
new equipment and plants, place Acroni, d. o. o., Jeseni-
ce at the same world quality level as other leading manu-
facturers of non-oriented electrical steels.
The losses of many grades with a low or medium
content of alloying elements have been reduced and in
some cases the permeability has been increased6. Lean
grades of non-oriented steel sheets have been developed.
A lower slab-re-heating temperature, the better defined
process of hot rolling and the higher final annealing
temperature are processing parameters that have been
used to improve the properties6.
3 CHARACTERISTICS AND MAGNETIC
PROPERTIES
The majority of reputable manufactures of electrical
machines will use fully or semi-processed silicon steel
with high quality. The principal quantity of interest for
soft magnetic materials is the power loss under
alternating current excitation (core loss) at a particular
operating frequency and at a particular maximum flux
density.
The electrical steel must satisfy several requirements,
with priorities that depend on the specific application,
such as high magnetic permeability, low hysteresis
losses, the anisotropy of the losses as well the ease of
cutting the laminations to shape. Various cutting
processes are applied, such as mechanical and laser
cutting. Mechanical cutting has been widely used in
industry due to its low cost. The magnetic properties of
the region at the edge are degraded after cutting. During
laser cutting, rapid heating and cooling cause thermal
stresses, which are also considered harmful for the
magnetic properties. On the other hand, by laser cutting,
the high temperatures may cause a grain growth near the
cut edge, which is beneficial for the magnetic
properties7.
Steel cuts-laminations are then built into the cores.
The laminations within the cores are physically rotated
relative to one another in order to equalize both the
reluctance of the flux paths within the material and the
variations in the thickness across its width. The designs
are usually optimized to utilize to the fullest the
magnetic and electric loading of the active materials,
copper and steel. This generally means that the steel is
pushed close to magnetic saturation and the copper/
insulation system is working close to its thermal limit8.
Electrical steel sheets are usually coated to increase the
electrical resistance between the laminations, to provide
resistance to corrosion or rust, and to act as a lubricant
during the cutting. There are various coatings, organic
and inorganic, and the coating used depends on the
application of the steel.
The magnetizing properties required for an NO
electrical steel sheet are achieved through measures such
as the purification of steel and the control of alloying
elements, the grain orientation and the grain size. When
the content of an alloying element such as Si is
increased, the electric resistance increases, the eddy-
current intensity in the steel sheet is decreased and as
result, the iron loss is reduced. However, the saturation
magnetic flux density is also reduced at the same time.
Thus, it is necessary to control the iron loss and the
saturation magnetic flux density in a well-balanced
manner.
Another factor to be considered is the influence of
magnetic domains. During an in-situ observation of the
magnetic domain structures a discontinuous movement
of irregularly distributed domain walls was observed, as
was the pinning of these domain walls by small oxide
precipitates and their strain fields. It was considered that
one of the causes of the domain-wall pinning is a
reduction in the magnetostatic energy inside the pre-
cipitates. The domain walls were observed to move in
D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
318 Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325
Figure 1: Coils of cold-rolled, non-oriented electrical steel5 (Acroni,
d. o. o., Jesenice)
Slika 1: Kolobarji hladno valjane neorientirane elektroplo~evine5
(Acroni, d. o. o., Jesenice)
curved lines around the precipitates in non-oriented steel
sheet.9,10
The domain width increases with the increase of the
grain size, resulting in an increase of the eddy-current
loss. As a result, a critical grain size exists to decrease
the iron loss. The relationship between the grain size and
the domain-wall width is given as follows:11
d3/4 = lg (/K1)
(/1.32)
d … the grain size
 … the domain-wall energy in a unit domain area
K1 … the magnetocrystalline anisotropy constant
d … the domain width.
If any strain or stress remains in the steel sheet, its
magnetic domain structure becomes complicated, the
magnetizing properties are deteriorated and the iron
saturation loss is increased. Thickness also significantly
influences the iron loss. The thinner a steel sheet is, the
more the eddy-current intensity is decreased.12 The
eddy-current losses are proportional to the square of the
frequency and the thickness of the sheet (the current
loops appear in the sheet section perpendicular to the
magnetic flux, and create a counter-field which opposes
the induction of the induction field). But when the steel
sheet is too thin, the iron loss increases rather than
decreases.12
Magnetic polarization and specific total loss are
measured by applying the Epstein method (EURO-
NORM 118, IEC 404-2). Sometimes a single-sheet tester
(IEC 404-3) may be used as an alternative4.
4 APPLICATIONS
The laminations form the laminated cores of trans-
formers or the stator and rotor parts of electric motors.
There is a wide range of equipment in which non-
oriented electrical steel sheets are incorporated, from the
simplest domestic appliances to hybrid electric vehicles.
Modern technologies were developed for the hybrid
electric vehicle, which is driven by an internal
combustion engine and an electric motor to lower the
fuel consumption and decrease the emission of exhaust
gases. Electrical steel sheets used for the core of the
traction motors of hybrid electric vehicles (HEV) and
electric vehicles (EV) affect the performance of
HEV/EV. The demand for smaller, lighter, more
powerful and more efficient motors is the driving force
for the development of electrical steel sheets.
5 METALLURGY
Iron-silicon alloys used for magnetic applications are
known as silicon steels. The production process of these
steel for NO electrical steel sheets and its chemical
composition are left to the discretion of the manufac-
turer. The iron–silicon binary phase diagram is shown in
Figure 2.
The a  	  transformation temperature is increased
and that of the  	  transformation is lowered until the
two meet at about 2.5 % Si, forming a closed "gamma
loop". As a result, an alloy containing more than about
2.5 % Si is body-centered cubic at all temperatures up to
the melting point. The a solid solution of the silicon in
iron is often called silicon iron. The presence of carbon
widens the ( + ) region, and only 0.07 % C shifts the
nose of the gamma loop to about 6 % Si14. In practice,
the carbon content in electrical steels is much lower, less
than 0.01 % C. A typical microstructure of a fully
processed, non-oriented, electrical steel sheet is shown in
Figure 3.
The addition of silicon to iron has the following
effects on its (magnetic) properties:14
1. The electrical resistivity is increased, the eddy-
currents are diminished and the losses are lowered.
2. The magnetocrystalline anisotropy decreases, causing
an increase in the permeability.
Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325 319
D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
Figure 2: Binary phase diagram of Fe-Si13
Slika 2: Binarni fazni diagram Fe-Si13
Figure 3: Polygonal grains of ferrite in the microstructure of the fully
processed, non-oriented, electrical steel sheet (LM, mag. 100-times;
etchant: Nital)15
Slika 3: Poligonalna feritna zrna v mikrostrukturi izdelane neorienti-
rane elektroplo~evine (SM, 100-kratna pove~ava; jedkalo: Nital)15
3. The magnetostriction decreases, leading to smaller
dimensional changes with magnetization and demag-
netization, and to a lower stress-sensitivity of the
magnetic properties.
4. The saturation induction decreases.
5. When the Si content is higher than 3%, the brittleness
of the steel is increased and the cold deformability is
significantly impaired.
Nowadays, other alloying elements instead of, or in
addition to, silicon are widely used. Among them the
most important is aluminium, which affects the magnetic
properties of iron similarly as silicon does. For non-
oriented electrical steels with aluminium addition the
sum of contents of both base elements (Si + 2Al) is up to
4 % 1.
Al and Mn form the non-metallic inclusions AlN and
MnS in the steel; however, impurity elements like Cu, Ti,
Se, Cr, Zr etc. can also form inclusions and thus
influence both the texture development and the magnetic
properties.
The fabrication route for non-oriented electrical
steels includes:
– for fully processed, non-grain-oriented electrical
steels the process route is: steel making, casting, hot
rolling, pickling with or without annealing, cold
rolling in one or two steps with an intermediate
annealing, final annealing and coating.
– for semi-processed material grades, a temper rolling
follows the annealing. The final annealing of the
stamped parts takes place at the customer’s site.
With the final annealing of the semi-processed
materials decarburization, surface oxidation and/or grain
growth are achieved and the required magnetic properties
are obtained1.
Depending on the alloy type (e.g., silicon content)
and the fabrication process, the hot rolling is carried out
as austenitic, two-phase, mixed or ferritic rolling.
Typically, a hot band with a thickness in the range
2.0–3.0 mm is used for the production of NO-material
grades. With respect to the steel’s final thickness of 0.35
mm, the total cold-rolling deformation is fixed to values
smaller than 85 % 1.The hot and cold rolling in
combination with the thermal treatment (annealings) and
the variation of the composition of the alloy are the
processing variables for achieving the required magnetic
and other physical properties of the silicon steel sheets.
The decarburization of the cold-rolled steel sheets is
a very important processing step because the texture
development and the magnetic properties are strongly
dependent on the carbon content. Iron carbides can
precipitate and degrade the magnetic properties by
interfering with the magnetic domain-wall motion. The
slow precipitation of the carbides during service is called
"magnetic aging" and can cause a substantial increase in
the core losses14. This magnetic aging anisotropy can be
associated with the crystallographic and morphological
characteristics of cementite (Fe3C) and -carbide (Fe2.4C)
precipitates formed during the aging treatment, taking
into account the texture developed in the steel.
The decarburization is performed by annealing in a
gas mixture of hydrogen and water vapor H2-H2O with a
controlled partial pressure ratio of water vapor and
hydrogen p(H2O)/p(H2) in the temperature range from
700 °C to 900 °C. The decarburization process of steel
consists of:16
1. Diffusion of carbon to the steel surface
2. Transport of water vapor to the steel surface and
equilibration at the phase boundary steel-gas mixture
3. Dissociation of water vapor molecules into hydrogen
and oxygen and adsorption on the steel surface
4. Oxidation of carbon
5. Oxidation of iron and alloying elements
The decarburization proceeds predominantly accord-
ing to the reaction:
[C]Fe + H2O(g) = CO(g) + H2(g).
The reaction:
[C]Fe + 2 H2(g) = CH4(g)
can be neglected at a p(H2O)/p(H2) greater than 0.01 17.
The thermodynamical calculations of equilibrium of
complex reactions for various furnace temperatures and
gas compositions have shown the conditions (gas-com-
position, temperature) for the formation of an oxide-
scale on the surface of non-oriented electrical sheet steel
during the decarburisation and thermal processing in
industrial continuous furnaces18.
While carbon is oxidized to the gases CO and CO2,
the steel surface is continually oxidized to a scale layer,
which is influenced by alloying elements, affecting the
oxidation process of iron. A typical oxide layer on the
steel surface of an Fe-Si-Al alloy is designated by the
arrow in Figure 419.
The decarburization annealing of electrical steels has
a significant effect on the final magnetic properties. The
process coarsens the grain size, removes the harmful
D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
320 Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325
Figure 4: Oxide layer on the steel surface of an Fe-Si-Al alloy for
non-oriented electrical steel after decarburization annealing (LM,
non-etched)15
Slika 4: Oksidna plast na povr{ini zlitine Fe-Si-Al za neorientirano
elektroplo~evino po `arjenju za razoglji~enje (SM, nejedkano)15
effect of carbon, but it can produce a strong texture in the
sheet. Namely, during the decarburization annealing of
cold-rolled sheets also the recrystallization processes
take place that may lead to the generation of texture
components unfavorable to the magnetization and
thereby adversely affect the magnetic properties of the
material. The temperature profile during the decarburi-
zation process due to differences in the heating rate can
lead to significant changes in the mechanism of
grain-boundary motion. Some elements, e.g., antimony,
decrease the solubility of the carbon in ferrite, promote
the precipitation of carbides and decrease the decarbu-
rization kinetics20,21. For these reasons the decarburi-
zation process requires a judicious optimization, with the
aim to achieve a favorable recrystallization and
grain-growth process.
6 TEXTURE
Non-oriented electrical steels have been among the
steel products that benefit most from texture optimi-
zation for the improvement of magnetic properties;
however, the focus of processing technology has largely
been on the control of grain size. Grain-size optimization
has been achieved by controlling the chemical com-
positions and optimizing the processing variables during
each processing step.
In contrast, the control of texture has received little
attention; hence, there is an unexplored possibility of
improving the magnetic properties of non-oriented steels
through texture control. A combination of metallo-
graphic and texture analyses with the measurement of
magnetic properties on annealed specimens allows the
most important microstructure and texture evolution
stages to be distinguished.22
The texture is a population of crystallographic
orientations whose individual components are linked to
their location within the microstructure23.
The ideal texture for a non-oriented silicon steel is a
random cube texture (001)[uv0], where each grain has
the <100> plane in the sheet plane, and the properties are
nearly isotropic. However, no industrial process has so
far been developed to produce this ideal texture
commercially. Texture improvement has been achieved
mainly by reducing the volume fraction of the [111]IIND
fiber, which is the main recrystallization texture
component of a-iron, and increasing the volume fraction
of texture components belonging to [001]IIRD and
[001]IIND fibers24.
The experimental technique used nowadays to
determine crystallographic information in non-oriented
electrical steel sheets and other materials as well is
EBSD (Electron Backscatter Diffraction) in a scanning
electron microscope (SEM).
The forerunner of EBSD was first reported in the
1930s as observations of high-angle Kikuchi patterns.
The biggest step forward, which was to result in the
emergence of EBSD as a sophisticated experimental
tool, occurred when diffraction patterns could be viewed
live by video detection and indexed on-line. Nowadays,
patterns from any crystal system can, in principle, be
indexed automatically. A very exciting EBSD output is
an "orientation map", which is a quantitative depiction of
the microstructure in terms of its orientation consti-
tuents23.
In EBSD a stationary electron beam strikes a tilted
crystalline sample and the diffracted electrons form a
pattern on a fluorescent screen. This pattern is charac-
teristic of the crystal structure and the orientation of the
sample region from which it was generated. The
diffraction pattern can be used to measure the crystal
orientation, measure grain-boundary misorientations, etc.
When the beam is scanned in a grid across a poly-
crystalline sample and the crystal orientation is measured
at each point, the resulting map will reveal the consti-
tuent grain morphology, orientations, and boundaries25.
In Figure 5 an EBSD orientation map of the micro-
structure of a non-oriented electrical steel sheet is shown.
The recrystallization texture is determined by both
the orientation of nuclei in the deformed matrices and
the growth rate of these nuclei into the deformed matrix.
Two main theories of recrystallization texture have been
currently accepted after a controversy lasting for over 50
years. The first one, oriented nucleation theory, assumes
that nuclei of specific orientations are faster in forming
than those of other orientations, and consequently
determine the recrystallization texture. The second one,
oriented growth theory, claims that there exist specific
rotation relationships with rapid grain-boundary
migration.
It was shown26 for a steel with 2 % Si that the
formation of recrystallization texture is explained by
D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325 321
Figure 5: EBSD Orientation map of the microstructure of non-oriented electrical steel sheet (using coloring)15
Slika 5: EBSD-prikaz orientacije feritnih zrn mikrostrukture neorientirane elektroplo~evine (z uporabo barvne lestvice)15
oriented nucleation. Most nuclei have a high misorien-
tation angle of 25–55° with the surrounding deformed
matrices. New Goss grains are mainly nucleated within
shear bands in the deformed {111}p, {111}n and {112}n
grains, and the number of shear bands decreases in the
same order. The nucleation of new cube grains also takes
place within the shear bands. New {111}p grains are
nucleated within the deformed {111}n grains and new
{111}n grains originated in the deformed {111}p grains.
The influence of the applied thermo-mechanical treat-
ments on microstructure progress in the non-oriented
steels was also studied. A columnar microstructure can
be obtained after combining the temper rolling and
appropriate annealing conditions. It was confirmed that
the obtained columnar microstructure possesses pro-
nounced cube-texture components.27
The evolution of the texture during the processing of
silicon steels for non-oriented electrical steel sheets and
the influence of small additions of the surface-active
element antimony on the recrystallization and texture
formation of silicon steels have been studied28-31. The
positive effect of antimony addition to the silicon steel
was reflected in a greater remanent induction and a lower
coercive force, which should lead to a smaller area of the
demagnetization loop and to smaller inductive-energy
losses for the NO electrical steel sheet31.
Later, the kinetics of the surface segregation of
antimony in silicon steel was investigated32 and the
segregation of antimony at the grain boundaries of a-iron
was also quantitatively determined33. It was confirmed
that the positive effect of antimony on the recry-
stallization behavior and on the texture of the silicon
steel is related to the surface segregation of antimony34.
A similar segregation propensity was observed for tin.
During the recrystallization annealing tin segregated to
the surface and decreased the surface energy selectively
and also selectively increased the mobility of some of the
grain boundaries. By alloying the silicon steel with
0.05 % Sn, a positive effect on the texture development
was achieved35,36.
The surface segregation of selenium was also studied.
It was found that this segregation critically affected the
reconstruction of the (110) surface of the grains in the
FeSi alloy, resulting in the formation of (100) planes at
850 °C (Figure 6).37
Since scrap steel is used in the industrial production
of non-oriented electrical steel the content of an
impurity, copper, is constantly increasing in steel. The
surface segregation of copper in silicon steels was
measured in-situ (Figure 7) in the analysis chamber of
an Auger spectrometer38,39. It was found that the intensity
of the surface segregation of copper increased with
increasing annealing temperature (Figure 8) and the
process of the surface segregation of copper during the
annealing of Fe-Si-Al alloys was described by the
dynamic equilibrium: 38,39
Cu(dissolved) = Cu(segregated) 	 Cu(desorbed)
Using EBSD it was possible to determine that the
mictrotextures of silicon steel sheets containing copper
D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
322 Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325
Figure 8: The surface segregation of Cu in Fe-Si-Al alloys38,39
Slika 8: Povr{inska segregacija Cu v zlitini Fe-Si-Al38,39
Figure 6: SE image of the surface reconstruction during the heating of
a selenium-doped FeSi alloy37
Slika 6: SE-posnetek rekonstruirane povr{ine med `arjenjem zlitine
FeSi, legirane s selenom37
Figure 7: Resistive heating of a steel sample in the ultra-high vacuum
of an Auger spectrometer for in-situ studies of surface-segregation
phenomena; Institute of Metals and Technology, Ljubljana.38
Slika 7: Elektrouporovno gretje vzorca v ultravisokem vakuumu
Augerjevega spektrometra, kar omogo~a in-situ analizo procesov
segregacije; In{titut za kovinske materiale in tehnologije, Ljubljana.38
had fewer crystal grains oriented with the easy axis of
magnetization and that grains with hard orientations
were more numerous (Figure 9)40.
The required texture of non-oriented electrical steel
sheets can only be obtained by controlling the content of
alloying and the trace elements40–44, and the dispersion of
precipitates and inclusions45,46, which all influence the
the recrystallization and grain growth processes.
7 NON-METALLIC INCLUSIONS AND
PRECIPITATES
The non-metallic inclusions in steels can be classified
as primary inclusions, formed during the refining stage,
and secondary inclusions, precipitated during the solidi-
fication and afterwards. Depending on the generation
source, non-metallic inclusions can be endogenous and
exogenous. Numerous studies focusing on the correlation
between the chemical composition and the morphology
of the inclusions and precipitates in electrical steels have
been performed over the past 15 years.44,47–59
The size of the non-metallic inclusions is normally
around 1 μm or more. The inter-inclusion distance is
much larger (micrometer scale). On the other hand, the
size of the precipitates is within the nanometer range.
The inter-particle distance can also be very small. These
small precipitates precipitating in the solid phase
deteriorate the magnetic properties of electrical steels by
pinning the motion of the domain walls.9,10
The precipitation reactions in alloys are thermally
activated atomic movements and are induced by the
change of the temperature of an alloy that has a fixed
bulk composition. From a metastable supersaturated
solid solution, stable or metastable precipitates are
formed, resulting in a more stable solid solution with a
composition closer to the equilibrium.
Oxygen, sulfur, and nitrogen are chemical elements
that have a decreasing solubility in iron with a decreas-
ing temperature. To prevent any deleterious precipitation
in electrical steels their content should be rigorously
controlled.44 One very harmful feature is the effect of the
sulfides, carbides or nitrides in the size range 10–400
nm, which is stronger with a greater density of particles
per unit volume57.
The stability of precipitates in silicon steels depends
on the temperature (Figure 10) and it is the greatest for
AlN.57
The precipitation of AlN and MnS is used to control
the texture development in electrical steel sheets. The
presence of these particles plays an important role in the
formation of the Goss texture. A classical method used
by metallurgists is to add to the alloy a solute element
with a low solubility, which precipitates as second-phase
particles that are able to pin the grain boundaries at low
temperature and allow grain growth at high temperature.
A typical AlN inclusion formed during the solidifi-
cation of steel is shown in Figure 11.
A prerequisite for the formation of a high density of
grains with the Goss texture in NO electrical steels
sheets is the inhibition of recrystallized grain growth up
D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325 323
Figure 11: SE image of AlN in a non-oriented electrical steel sheet15
Slika 11: SE-posnetek AlN v neorientirani elektroplo~evini15
Figure 10: Precipitates in silicon steel57
Slika 10: Precipitati v silicijevem jeklu57
Figure 9: Inverse pole figure of an industrial, non-oriented, electrical
steel containing 0.60 % Cu38
Slika 9: Inverzna polova figura industrijske neorientirane elektroplo-
~evine z 0,60 % Cu38
to a temperature above 1050 °C, when the conditions for
the texture formation are achieved by annealing in dry
hydrogen. AlN and MnS inclusions inhibit the grain
growth. The solubility products are therefore very
important for the texture formation. The solubility of the
complex sulfide (MnxFe1–x)S in a 3 % Si steel in the
temperature range from 1100 °C to 1300 °C was calcu-
lated53. This solubility was in good agreement with the
analyzed microstructures and precipitates. Some results
show50 that the particles of AlN and MnS in non-oriented
electrical steel grades have grown, to some extent, in the
soaking stage. The temperature range from 1000 °C to
1200 °C is appropriate for reheating before hot rolling
for the majority of steels.60 It is reported that the effect of
particles AlN and MnS precipitated during and after the
hot rolling is much stronger. In non-metallic inclusions
in the selenium-containing, non-oriented, electrical steel
sheet, both copper-selenides and complex copper-
selenide inclusions were found. The complex selenides
were found to grow on the nitride, oxide or oxysulfide
particles in the steel51.
Using a statistical multivariate analysis on a data set
of 409 coils of non-oriented electrical steel sheet it was
found that the titanium content has a strong and negative
effect on the core losses, much greater than other
elements. The trend of its influence is similar to that of
copper61 (Figure 12).
Titanium is a non-magnetic element and as such
diminishes the saturation magnetization. For this reason
the negative effect of titanium shown in Figure 12 could
be explained by the presence of titanium precipitates of
the nitride, carbide, and carbonitride type. This is similar
to the effect of copper, which has a much stronger
negative effect on the magnetic properties in the form of
precipitates than as present in the solid solution.
To ensure the improved magnetic properties of NO
electrical steel sheets the content of the impurity
elements should be minimized44. Namely, the process of
grain growth in the primary recrystallized matrix mainly
depends on the number and the dispersion of second-
phase particles46. Particles that are precipitated from a
supersaturated solid solution can have a negative effect
on the secondary recrystallization and also on the texture
formation58,59. Nanoscale precipitates play an important
role since they can hinder the process of magnetization
by pinning the wall movements of the magnetic
domains62.
Generally, for good soft-magnetic properties, i.e., low
losses and a good permeability, electrical steels should
have as few precipitates as practically and economically
possible.
Recently, a new generation of high-permeability,
non-oriented grades was developed, based on an
improved crystallographic texture and purity. Due to the
beneficial effects of texture and purity on the core losses,
the new grades could be produced with a lower alloy
content (Si+Al), with positive effects on the mechanical
properties, the thermal conductivity, the saturation
polarization and the magnetic permeability24.
8 CONCLUSIONS
The metallurgy of silicon steels for non-oriented
electricals steel sheets is very complex. Numerous
precautions must be taken during the manufacturing
process. Nevertheless, the processes are now well
controlled and little further improvement can be
expected from simple compositional modifications.
Future efforts must be focused on controlling the
residual elements in the steel melts, optimizing the hot
and cold rolling and optimizing the crystallographic
texture development in order to enhance the performance
of the finished product, since being environmentally
friendly is one of the essential requirements for the
future.
The market for electrical steels is large, since there is
a wide range of equipment, from the simplest domestic
appliances to heavy electrical engineering applications.
The introduction of high-permeability products and
particularly the development of hybrid electric vehicles
will be a major source of the expansion of non-oriented
electrical steel sheets in the future.
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D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
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D. STEINER PETROVI^: NON-ORIENTED ELECTRICAL STEEL SHEETS
Materiali in tehnologije / Materials and technology 44 (2010) 6, 317–325 325

M. TORKAR: EFFECT OF TRACE AND RESIDUAL ELEMENTS ON THE HOT BRITTLENESS ...
EFFECT OF TRACE AND RESIDUAL ELEMENTS ON
THE HOT BRITTLENESS, HOT SHORTNESS AND
PROPERTIES OF 0.15–0.3 % C Al-KILLED STEELS WITH
A SOLIDIFICATION MICROSTRUCTURE
U^INEK ELEMENTOV V SLEDOVIH IN PREOSTALIH
ELEMENTOV NA KRHKOST IN POKLJIVOST V VRO^EM JEKLA
Z 0,15 – 0,3 % C, POMIRJENEGA Z Al IN S STRJEVALNO
MIKROSTRUKTURO
Matja` Torkar
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana
matjaz.torkar@imt.si
Prejem rokopisa – received: 2010-09-09; sprejem za objavo – accepted for publication: 2010-11-11
The formation of cracks in steels with a solidification microstructure is connected with either the precipitation and segregation
of residual elements during solidification and/or with the accumulation of some residual elements at the scale–metal interface
during the soaking of the blocks. Hot-brittleness is connected with the change in the solubility of the trace elements in the solid
solution of austenite, with segregations of the surface-active elements to the grain boundaries, with the precipitation of small
particles at the grain boundaries during solidification and cooling to the hot working temperature, and with the effect of some
elements on the solidification microstructure. In this article an overview of the hot brittleness of steels cast in the peritectic
range is presented.
Key words: solidification, hot-brittleness, hot-shortness, segregation, precipitation, residuals, selective oxidation, grain-
boundary penetration, hot-cracks
Nastanek razpok v jeklih s strjevalno mikrostrukturo ima dva vzroka: izlo~anje in izcejanje elementov v sledovih med
strjevanjem/ali bogatenjem vsebnosti nekaterih oligoelementov na meji {kaja-kovina med ogrevanjem blokov. Vro~a krhkost je
povezana s spreminjanjem topnosti oligoelementov v trdni raztopini v avstenitu, z izcejanjem povr{insko aktivnih elementov po
mejah zrn, z izlo~anjem drobnih delcev po mejah zrn med strjevanjem in ohlajanjem na temperaturo vro~e predelave ter z
u~inkom nekaterih elementov na strjevalno mikrostrukturo. V tem ~lanku je prikazan pregled vro~e krhkosti litega jekla v
obmo~ju peritektika.
Klju~ne besede: strjevanje, krhkost v vro~em, pokljivost v vro~em, izcejanje, izlo~anje, oligoelementi, selektivna oksidacija,
penetracija po mejah zrn, razpoke v vro~em
1 INTRODUCTION
The increasing use of steel scrap in electric arc
furnaces means a constant increase in the quantity of
residual elements in steels. These residual elements (Cu,
Ni, As, Pb, Sn, Sb, Mo, Cr, etc.) are defined as elements
that are not added to steel and which cannot be removed
by current metallurgical processes. The effect of some
residual elements is connected with their presence in the
solid solution, such as Mo, Cr, Ni, and Cu, while others
segregate at the interfaces (surfaces and grain boun-
daries), such as Sn, As, and Sb1, and some effect the
solidification microstructure, as Al and N. Some of these
elements can be removed by appropriate processing,
while some cannot be removed from the molten steel and
induce hot brittleness and hot shortness in the steels. The
brittleness is found mostly for steel with carbon in the
peritectic range of 0.07 % to 0.18 % C. During the
solidification of steel the trace elements and other
impurities concentrate in the last melt and are pushed by
the solidification front towards the interdendritic regions.
There they may facilitate the formation of cracks or
cause brittleness of the steel when the steel is exposed to
shrinkage stress. Some of the residual elements also
influence the solidification mode, while some impurities
remain in solid solution during rapid solidification and
can precipitate from an oversaturated solid solution on
the grain boundaries during the cooling of steel after hot
working or heat treatment and induce the hot brittleness.
Manganese, aluminium and carbon influence the
interparticle distance and the size of the sulphide parti-
cles that occur in steel during cooling from the tempera-
ture of solidification. An increasing content of manga-
nese increases the size and the interparticle distances,
which are lower in steels containing aluminium. It seems
that sulphur originates from the segregation during
solidification (Figure 1), but to draw a conclusion about
where the intercrystalline sulphide really originates from
is not yet possible2,3.
The increasing aluminium content promotes a
columnar solidification structure in steel. In steel with
high aluminium content, coarse columnar grains at the
surface of the billets facilitate the propagation of
Materiali in tehnologije / Materials and technology 44 (2010) 6, 327–333 327
UDK 669.14:620.1/.2 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 44(6)327(2010)
intergranular cracks in length and depth (Figure 2).
After cooling to ambient temperature and reheating, the
structure was identical at all levels of aluminium content
and consisted of regular polygonal grains of smaller size.
The effect of aluminium should be attributed to a process
or a reaction connected with homogeneous nucleation at
the start of the solidification3. The hot shortness of steels
with carbon in the peritectic range is observed mostly in
steels with a high content of nitrogen and refined in an
electrical arc furnace. This suggests a synergistic effect
of aluminium and nitrogen on the solidification struc-
ture4. With the addition of aluminium the undercooling
of the liquid steel is decreased. This reduces the
possibility of the formation of a sufficient globular-
grains layer near the surface of the billet. If a stronger
nitride-forming element, such as titanium, is introduced
into the steel, the hot shortness is greatly decreased by an
equal content of aluminium and nitrogen. Besides that,
every process which modifies the solidification structure
of the steel, e.g., recrystallisation and the transformation
of austenite, improves the workability of the steel to a
M. TORKAR: EFFECT OF TRACE AND RESIDUAL ELEMENTS ON THE HOT BRITTLENESS ...
328 Materiali in tehnologije / Materials and technology 44 (2010) 6, 327–333
Figure 3: Concentration of copper and tin at the scale–metal interface
at various annealing temperatures. Steel contains 1 % Cu and 0.030 %
Sn5.
Slika 3: Koncentracija bakra in kositra na prehodu iz {kaje v kovino
pri razli~nih temperaturah `arjenja. Jeklo ima 1 % Cu in 0,030 % Sn5
Figure 1: Distribution of S and Mn near the solidification boundary.
a) 0.15 % C, 0.18 % Si, 0.33 % S (in mass fractions), specimen cast,
stripped and quenched; b) 0.17 % C, 0.20 % Si, 0.14 % Mn, 0.022 % S
(in mass fractions), specimen cast, stripped, held for 30 min at 1250
°C and quenched, i. p. is an intergranular precipitate; c) 0.17 % C,
0.02 % Si, 0.022 % S (in mass fractions), specimen cast, stripped, held
for 8 h at 1250 °C and quenched3
Slika 1: Razporeditev S in Mn blizu strjevalne meje: a) 0,15 % C,
0,18 % Si, 0,33 % S (mas. dele`i), vzorec ulit, izvle~en in ohlajen v
vodi; b) 0,17 % C, 0,20 % Si, 0,14 % Mn, 0,022 % S (mas. dele`i),
vzorec ulit, izvle~en, zadr`an 30 min na 1250 °C in ohlajen v vodi; i.
p. so interkristalni izlo~ki; c) 0,17 % C, 0,02 % Si, 0,022 % S (mas.
dele`i), vzorec ulit, izvle~en, zadr`an 8 h na 1250 °C in ohlajen v
vodi3
Figure 2: Influence of Al content on the thickness of the equiaxed
shell and the length and width of the columnar grains4
Slika 2: Vpliv vsebnosti Al na debelino plasti enakoosnih zrn ter
dol`ino in {irino stebrastih zrn4
great extent and a surface free of rolling defects is
obtained on the hot-rolled products.
Hot-shortness occurs during the hot working of steel.
Because of selective oxidation, Cu, Sn and other trace
elements concentrate in the solid and liquid state at the
scale–metal5 interface (Figure 3). Penetration with
grain-boundary diffusion along the grain boundaries
(Figure 4) weakens the cohesion of the grains and
induces surface cracks during the hot working6.
The general influences of residuals on the properties
of steel are given in Table 1.
Table 1: Effects of the increase in the amount of residual elements on
various properties of steel1
Property Cu Ni Cr Mo Sn Sb
Strength and hardness + + +,– + + +
Ductility – +,– +,– – –
Strain hardening, n – – – – –
Strain ratio, r +,– –
Impact resistance + + – –
Hardenability + + + + +
Weldability – – – –
Corrosion resistance + + + +
Temper embrittlement + +
Elongation – – –
(+) increase of, (–) decrease of
A large proportion of modern steel production is
based on the recycling of steel scrap and the use of only
an electric arc furnace (EAF) as a melting aggregate. All
the other processing occurs in separate ladle furnaces as
secondary or vacuum metallurgy with alloying,
rafination and achieving the proper casting temperature.
Advances in technology make it possible to remove the
sulphur and phosphorus from the steel much more
effectively than in the past. But the increased use of
scrap in an EAF increases the content of residuals like
copper and tin in the steel and causes problems with the
hot shortness and hot brittleness of the steels. For this
reason, the proper selection of the scrap is of great
importance.
The problems with the hot brittleness of steel were
intensively investigated in previous years. Numerous
studies and investigations have been performed since the
early 1900s, when steel’s propensity for hot cracking
was referred to as red-shortness.
The cause of more intensive investigations in the late
1950s and 1960s was the increase in the amount of steel
production from scrap that caused a larger amount of
copper and other residuals in the steels7 and again in the
late 1990s, for economic and environmental reasons8.
The beneficial effect of nickel on the reduction of the
detrimental effect of copper was recognized early. Hot
brittleness and hot shortness are still encountered and
attract more intensive investigations from time to time.
In this article the present state of knowledge with
regard to both phenomena is presented.
2 HOT BRITTLENESS OF AS-SOLIDIFIED
STEEL
The introduction of microprobe analyses (MPA) in
the late 1960s and 1970s produced new knowledge on
the hot- brittleness and hot-shortness of steel. The goal
of the investigations was to find a mechanism for these
phenomena. The intensive study of the solidification
processes was performed for the influence of aluminium
on the initial deformability of continuously cast
C-Mn-Si-N steel4 and revealed that aluminium influen-
ces the hot workability via its influence on the solidi-
fication structure. Besides that, the influence of
manganese, carbon and aluminium on sulphur solubility
during solidification was determined by direct micro-
probe analysis. As no gradient of sulphur was found in
the grain-boundary areas, it was concluded that inter-
granular particles of MnS do not originate from sulphur
in solid solution in the interior of the solidification grains
but from grain-boundary segregation during solidifi-
cation3. The intergranular brittleness due to sulphide
precipitates causes a severe hot shortness of as-cast
0.16 % C steel with increased contents of aluminium and
nitrogen cooled from solidification to the rolling tempe-
rature9.
Sulphide precipitates are one of the reasons for the
hot brittleness of steel with an as-solidified structure and
hot worked at a temperature below the austenite transfor-
mation9. An increasing content of manganese increases
the size and interparticle distances, which are lower in
steels containing aluminium10,11,12.
Based on the results of investigations and on
reference data, two hypotheses are consistent with the
data in the references and with the findings of
investigation as to why the increased aluminium content
reduces the susceptibility of steel to form equiaxed
crystals in a layer at the surface and in the core of the
billets. One suggests that aluminium widens the region
M. TORKAR: EFFECT OF TRACE AND RESIDUAL ELEMENTS ON THE HOT BRITTLENESS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 327–333 329
Figure 4: Concentration of Cu-rich phase below the scale after 4 h of
annealing at 1050 °C. Etched with Dickenson reagent6.
Slika 4: Zbiranje s Cu bogate faze pod {kajo med 4-urnim ogrevanjem
pri 1050 °C. Jedkano z Dickenson jedkalom6.
of existence of -ferrite and the second one suggests the
formation of associations of aluminium and nitrogen
atoms in the melt, decreasing the undercooling of the
melt, which decreases the thickness of the surface layer
of equiaxed small grains and exposes the coarse
columnar grains to deformation with boundaries perpen-
dicular to the direction of steel deformation.
The hypothesis that hot brittleness is associated with
the influence of aluminium and nitrogen on the forma-
tion and thickness of a globular layer at the surface of
steel, cooled from solidification to the rolling tempera-
ture,13 was confirmed with industrial rolling tests of
billets from the same charge but with different contents
of Al (Figure 5 and 6).
Many elements segregate to the surface or to the
grain boundaries14,15, like C, Si, N, S, P, Sn, As, with
them all being in competition. Tin segregates at the MnS
particles and retards the growth of the MnS particles.
This has been observed in Fe-3 % Si electrical steel
grades16.
The main reasons for the hot brittleness of steel with
an as-cast structure are the effects of elements on the
solidification structure and the segregation of surface-
active elements or the precipitation of particles on grain
boundaries.
3 HOT SHORTNESS OF SOAKED STEEL
BLOCKS
It is agreed that the behaviour of copper and tin
during the reheating for the hot working of steel is the
main reason for the appearance of hot shortness. It was
established17 that tin in the absence of copper is not
detrimental to surface cracks. Antimony is likely to be
nearly as detrimental as tin because it reduces the
solubility of copper in austenite18.
The most investigated was the effect of copper and of
its diffusion. Several studies have shown that at high
temperatures the diffusion of copper along grain
boundaries prevailed over volume diffusion. Also, more
detailed data on the time-dependency of the solubility of
copper in steel were established and the conclusion was
that the effect of copper diffusivity depended on the
temperature.19,20,21,22,23
Hot shortness and the resulting increased propensity
for cracking are commonly associated with the presence
of liquid copper and/or copper-enriched phases at the
scale–metal interface.24,25
Two stages are involved in the phenomenon: the
formation of copper-rich inserts and the penetration
along the grain boundaries during reheating and the
crack formation during hot working. The scaling rate
(temperature, time, and atmosphere) and the deformation
temperature are of prime importance for this pheno-
menon. A complete issue of ISIJ International 37 (1997)
3 is dedicated to the copper-iron system and to trials of
the processing modification for the lowering of hot-
shortness sensitivity.
Intensive studies of the metallic compounds in the
scale layer, below the scale and in the metal26,27,28
identified the phases known from the appropriate ternary
phase diagrams. The beneficial influence of nickel,
found in earlier investigations29, was re-evaluated. It was
established that to prevent hot-shortness, the content of
nickel should be half that of the copper content. The
beneficial effect of the soaking temperature range
1200–1300 °C, with or without the presence of nickel,
was also reconfirmed. The main reason for the absence
of cracks in this high-temperature region is the rapid
diffusion of the copper in the austenite. The diffusivity of
copper in iron is five times higher at 1200 °C than at
M. TORKAR: EFFECT OF TRACE AND RESIDUAL ELEMENTS ON THE HOT BRITTLENESS ...
330 Materiali in tehnologije / Materials and technology 44 (2010) 6, 327–333
Figure 6: Defects on transverse section of (26 × 70) mm bar rolled in
solidification heat. a) 0.004 % Al, b) 0.04 % Al4.
Slika 6: Napake na pre~nem prerezu valjane palice prereza (26 × 70)
mm, ki je bila izvaljana takoj po strjevanju: a) 0,004 % Al, b) 0,04 %
Al4.
Figure 5: Longitudinal section of cracks at the bent edge of a billet
cut. a) shallow non-propagating cracks with blunted tip (0.004 % Al),
b) crack propagates to much greater depth (0.04 % Al)4.
Slika 5: Vzdol`ni prerez razpok na upognjenem robu odreza gredice:
a) plitve razpoke z zaobljenim vrhom, ki ne napredujejo (0,004 % Al),
b) razpoka, ki je napredovala mnogo globlje (0,04 % Al)4.
1100 °C. The consequence of this increased diffusivity is
that despite the much higher oxidation rate at 1200 °C
than at 1100 °C, the amount of copper-enriched phase at
the steel–scale interface was smaller at 1200 °C,
compared with that at 1100 °C.
Several studies have focused on the influence of
copper and tin on the hot ductility in steels 26,27,28,29,30.
Hot-shortness cracks develop because of surface tensile
stresses and boundaries weakened by the grain-boundary
copper penetration. The experiments confirmed that the
critical temperature of 1100 °C was just above the
melting point of copper and that the temperature of 1200
°C was above the reported critical hot-shortness tempe-
rature27,28,31,32,33,34,40. Above 1100 °C the copper-rich
phase is in the molten state at the scale–metal interface.
More critical are the lower temperatures when the
diffusion of copper along the grain boundaries prevails.
The enrichment process is balanced by the dispersion of
copper below the scale and the diffusion of copper in
iron26. Thus, there are two competing mechanisms,
grain-boundary penetration and matrix diffusion. The
steel’s hot-working temperature when the hot-shortness
cracks appear, is the breaking point of these two
mechanisms.
Copper enrichment at the scale–metal interface
occurs due to the selective oxidation of the steel surface
and because copper is insoluble in the scale. The
oxidation process is a parabolic function of time, and
numerous experiments confirmed that the enrichment of
copper below the scale depends on the extent of the
scaling. The oxidation rate of the steel’s surface is
temperature dependent, and for this reason the
enrichment rate is both temperature dependent and
dependent on the content of copper in the steel. For a
given temperature, the higher the content of copper in the
steel, the higher will be the enrichment. As the tempe-
rature increases, both the oxidation and the enrichment
rate increase also. At first, the grain-boundary pene-
tration predominates up to the point where the diffusion
coefficient of the copper is sufficiently high to start with
a planar diffusion front. The diffusion coefficient
increases exponentially with temperature. Thus, the
enrichment, the dispersion and the grain-boundary
penetration are interconnected mechanisms with a tem-
perature window of hot-shortness. The grain-boundary
penetration is a diffusion process and it is dependent on
the enrichment itself and on the temperature. Without a
sufficient amount of copper accumulation below the
scale, there is also no significant copper concentration at
the boundaries and no problems occur with hot-short-
ness.
For a high copper content the effect of nickel in steel
was investigated. It was found that the addition of nickel
to the steel prevented or decreased the hot-shortness for
steel with a high copper content (Figure 7). The
recommended addition of nickel was the same as the
content of copper. However, new experiments27,28 showed
that the content of nickel is half the copper content in the
steel to prevent the hot-shortness. Other investigations8,41
showed that in the case of a value of Cu:Ni ratio close to
4, the solidus line from the copper–nickel binary phase
diagram was approximately 1160 °C. Nickel increases
the Cu-solubility in the austenite as well as favouring the
occlusion of Cu in the scale. The effect of nickel results
in the modification of the chemical composition of the
phases, with a higher melting temperature, formed on the
steel surface24,26,35,36 that lower the propensity for hot
cracks’ formation.
The copper content varies among steels and
accordingly variations in the enrichment rate occur. The
variation in temperature affects the diffusivity of the
copper in the steel, and thus affects its dispersion rate.
For steels with a significant copper content, there are two
critical oxidation times: the first occurs soon after the
formation of the oxide layer, due to copper enrichment
from the oxidised steel, and rapid copper penetration into
the grain boundaries, that weakens their coherence. By
increasing the temperature the scale becomes thicker, the
oxidation rate is decreased, while the bulk diffusion of
copper is faster and the surface of the steel is less
sensitive to cracking. This explains why the steel with
M. TORKAR: EFFECT OF TRACE AND RESIDUAL ELEMENTS ON THE HOT BRITTLENESS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 327–333 331
Figure 7: Propensity to hot shortness during a hot-bend test of as-cast
steel ^ 4320, dependent on the Cu:Ni ratio and the Sn content. a)
0.015 % Sn, b) 0.023 % Sn41
Slika 7: Vro~a krhkost pri vro~em upogibnem preizkusu litega jekla ^
4320 v odvisnosti od razmerja Cu/Ni in vsebnosti Sn: a) 0,015 % Sn,
b) 0,023 % Sn41
copper behaves better at 1200 °C than at 1100 °C in
terms of hot cracking. Also, a low temperature and a
smaller amount of copper available delay the penetration
into the grain boundaries. These two effects explain the
maximum cracking being delayed to longer oxidation
times.
Several tests were performed8 to evaluate the newer
processing conditions used in modern forge shops that
might allow the use of higher copper contents for forging
without fear of major surface cracking. A short induction
heating time to a temperature of 1200 °C reduces the
problems with selective oxidation and the hot-shortness
of steel. The forgeability of such steel is similar to steel
with a very low copper content.
Besides the residual content of copper in the steel,
sometimes copper is added to the steel because of the
beneficial influence on some specific properties of the
steel. What is well known is that the addition of copper
to structural steel increases its atmospheric corrosion
properties (in Corten steel, by up to 0.50 % Cu). The
addition of Cu is also known to contribute moderately to
austenite hardenability37. During hot working, in addition
to micro-alloying, Cu (mass fraction >0.8 %) contributes
to the retardation of recrystallisation and is usually added
in a larger amount to achieve the expansion of the
unrecrystallised region by the solute-drag effect35. A Cu
addition also improves the bainitic hardenability and, in
the presence of a suitable microalloying addition like Ti
or B, bainite may form even under air-cooling
conditions38. Although the addition of Cu does not
effectively contribute to the lowering of the austenite
transformation temperature; its addition suppresses the
pearlitic transformation by stabilising the austenite at the
grain boundaries and triple points of ferrite grains37.
It was established that the Cu-rich phase is
distributed homogeneously in the martensitic matrix of a
cutting tool36–39. The copper precipitates from the solid
solution in ferrite without a uniform orientation and
copper precipitates with a size is of approximately 10 nm
are formed. These Cu-rich precipitates have a lubricating
effect on the contact between the steel and the cutting
tool and promote the outflow of the cutting heat. The
effect of the lubricant and heat conductivity reduces the
cutting-tool wear, improves the machinability and
increases the service life of the cutting tool39.
However, in the exploitation of the beneficial effect
of copper on the steel’s properties it is necessary to
consider the eventual effect of hot shortness. Thus,
special care is demanded during the reheating process
and the hot working of steels with an increased copper
content.
4 CONCLUSIONS
The motivation for all the investigations of hot
brittleness and hot shortness is the need to find a solution
for the hot working of steels with a high residual-
elements content without surface cracking.
It is clear that the content of residual elements in
steel will increase with time due to the increasing
quantity of scrap recycling. Therefore, it is necessary to
identify and to quantify any deleterious effects in order
to keep these effects within acceptable limits.
The hot brittleness of as-solidified steel is a conse-
quence of the influence of aluminium and nitrogen on
the solidification structure, the presence of the segre-
gations of surface-active elements and the presence of
MnS particles at the grain boundaries and residuals’
enrichment at the scale–steel surface.
Residual elements, or at least some of them, may
affect the processing conditions, and the mechanical
properties of the end product. These residual elements
are more detrimental in all applications that require
low-carbon clean steels, extra-low-carbon clean steel,
ultra-low-carbon–interstitials-free clean steel (ELC, LC,
ULC-IF) than for the medium- and high-carbon-alloyed
steels.
Due to the influence of residual elements and
impurities on the hot-brittleness, hot-shortness and on
other properties of steels, plants using recycled scrap for
steel production need to be increasingly concerned with
scrap-quality selection.
From the results of numerous investigations it seems
that shorter heating times, controlled atmospheres in the
reheating furnaces, the application of protective coating
or higher temperatures for the hot working are the
options in any further investigations. Each option has
detrimental and beneficial effects that need to be
considered during the steel’s hot-processing technology.
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kovalno sposobnost povr{ine litega jekla, @elezarski zbornik, 13
(1979) 4, 161–170
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@elezarski zbornik, 15 (1981) 2, 61–68
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of Copper Addition on Mechanical Properties of 4Cr16Mo, Acta
Metall. Sin. (Engl. Lett.), 21 (2008) 1, 43–49
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mechanical processing on mechanical properties of copper bearing
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84-030, The influence of tin and copper on hot shortness of structural
steel with as cast microstructure Ljubljana, september 1984
M. TORKAR: EFFECT OF TRACE AND RESIDUAL ELEMENTS ON THE HOT BRITTLENESS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 327–333 333

B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
THE INFLUENCE OF POWDER SELECTION AND
PROCESS CONDITIONS ON THE CHARACTERISTICS
OF SOME OIL-RETAINING SINTERED BRONZE
BEARINGS
VPLIV IZBIRE PRAHOV IN POGOJEV IZDELAVE NA LASTNOSTI
SAMOMAZALNIH SINTRANIH LE@AJEV VRSTE Cu-Sn
Borivoj [u{tar{i~, Matja` Godec, Aleksandra Kocijan, ^rtomir Donik
Institute of Metals and Technology, Lepi pot 11, 10000 Ljubljana, Slovenia
borivoj.sustarsic@imt.si
Prejem rokopisa – received: 2010-03-23; sprejem za objavo – accepted for publication: 2010-07-26
A chemical, microstructural and mechanical characterisation of sintered bronze bearings from three different producers was
performed. The investigated bearings differ considerably in their final functional properties in spite of the fact that they all have
similar bulk chemical compositions and a similar basic process of their manufacture. Therefore, on the basis of the performed
characterization of these bronze bearings, the reasons for their good or bad functional properties were analysed and reported.
Keywords: oil-retaining self-lubricating sintered bronze bearings, chemistry, microstructure and mechanical properties, analysis
and comparison, the influence of process conditions
Kemijsko, mikrostrukturno in mehansko smo opredelili vzorce sintranih le`ajev treh razli~nih dobaviteljev. ^eprav imajo vsi
trije le`aji podobno povpre~no kemijsko sestavo in naj bi bili izdelani po enaki konvencionalni sinter tehnologiji, se med seboj
mo~no razlikujejo po svojih kon~nih funkcionalnih lastnostih. Zato na osnovi izvedene karakterizacije le`ajev poizku{amo
analizirati vzroke za dobre oz. slab{e funkcionalne lastnosti posameznih le`ajev.
Klju~ne besede: naoljeni samomazalni le`aji, kemijska, mikrostrukturna in mehanska karakterizacija, analiza in primerjava
le`ajev, vpliv procesnih parametrov
1 INTRODUCTION
Porous, oil-retaining bearings are both the best
known and the most widely used bearing type1. They are
usually produced by conventional sintering technology,
i.e., the uniaxial automatic die compaction (ADC) of the
powder mixture and sintering in a protective atmosphere.
The essential element of this technology is to ensure the
appropriate open porosity that is able to retain and,
during hydrodynamic loading of bearing, supply of
suitable oil content into the gap between the shaft and
the bearing’s sliding surfaces. The porosity of this type
of bearings is between 15 % and 25 %, i.e., the so-called
open porosity. This means that the pores are mutually
interconnected. The appropriate open porosity of the
bearing enables its full impregnation. The open pores are
infiltrated to at least 90 %. The industrial impregnation
of the bearings with oil is performed in vacuum
chambers. First, the air is evacuated from the pores and
then the bearings are dipped into the special oil. The
average size of the pores and their size distribution
mainly depend on the powder used, as well as the
compaction and sintering conditions. Standard, sintered,
tin-bronze bearings have the mass fraction of Sn
approximately 10 %. The German standard DIN 30910,
Part 3, gives the composition and the properties2. In
sintering practice it is possible to use a powder mixture
of elemental powders (Cu, Sn etc.), as well as alloyed
powder with a homogeneous final chemical composition
of each of the powder particles (in our case Cu-10 %
Sn). In some cases, pre-alloyed powder, i.e., only a
partial alloying of basic Cu particles, can also be
used1,3,4. Generally, in all cases the powder particles have
an irregular shape. Atomized powders are more globular,
but pure elemental Cu powder produced by electrolysis
has a dendritic or needle-like shape. Cu powder can also
be produced by the reduction of ground copper oxide in
a continuous belt furnace with hydrogen, dissociated
ammonia or an endothermic atmosphere. Low-melting-
point tin powder, usually produced by air atomization,
has regularly (i.e., spherically) shaped particles5. The
selected powder, compaction pressure and sintering
conditions define the final functional properties of the
bearing.
Sintered, self-lubricating, oil-retaining Cu-10 % Sn
bearings from three different producers were investigated
and analysed in this study. The bearings exhibit a large
difference in quality and functionality. The bearing
designated as bearing A, built in motors for liquid flow
regulation and the opening of valves, endures under the
standard loading test conditions for more than 150 000
cycles (open/close). However, the bearings designated as
B and C can cope with much fewer (less than 25 000
cycles). Figure 1a shows the bearing C after just 15 000
Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342 335
UDK 621.822: 620.1/.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)335(2010)
cycles of operation. It has been found eccentricity and
visible tracks of wear debris. The reasons for such large
differences in the bearings’ behavior and life time were
therefore established. A complete analysis and compa-
rison of the bearings were performed. The results of the
analyses, as well as the findings and possible reasons for
the differences in the bearings’ behavior are given below.
2 EXPERIMENTAL
The present sintered, self-lubricating bearings should
be manufactured in accordance with the German
standard2 in quality Sint B-50. This declares the
chemical composition (Cu with 9 % to 11 % Sn), the
bearings’ porosity of 20 ± 2.5 %, the sintered density
from 6.8 g/cm3 to 7.2 g/cm3 and a hardness larger than
35 HB.
The average bulk chemical composition of the
investigated bearings was determined with Ion Coupled
Plasma – Atom Emission Spectroscopy (ICP-AES),
Perkin Elmer 2380. The investigated bearings were
delivered in the oil-impregnated condition. Therefore,
the samples of the bearings were de-oiled and then
dissolved in a concentrated solution of gliceregia. The
formed solution was then diluted, analysed and
compared with standard solutions.
The thermal removal of the oil has to be avoided
because it can change the composition and properties of
the bearings (oxidation, the low melting point of Sn).
The standard ISO 27386 recommends the removal of the
oil from the test pieces by solvent extraction with
sequential Soxhlet Extraction. This is a very difficult and
time-consuming process. For practical control purposes
other methods may be used. In our case, de-oiling was
performed with cyclic cooking in leach, ultrasound
cleaning, drying at 105 °C and weighing of the samples.
The oil density of the sintered bearings was deter-
mined picnometrically6 in distilled water at room tempe-
rature. The method is based on Archimedes’ principal.
Two bearings from each producer were measured.
The Vickers hardness, with a loading of 5 kg (HV5),
of the bearings was measured, initially, as proposed by
the German standard7. However, the indentation depth
was too large and, finally, the loading was decreased to
1 kg (HV1).
For the microstructural investigations under a light
(LM) and a scanning electron microscope (SEM), the
samples were cut into halves and put into the bakelite in
two directions with respect to the compaction direction
(parallel/perpendicular; Figure 1b). A standard metallo-
graphic preparation (grinding and polishing) of the
samples was then performed. Porosity and inclusions are
observed on the polished surface, but the microstructural
constituents are observed on the etched surface. In our
case, an alcohol suspension of ferric chloride and hydro-
chloric acid (5 g FeCl3+10 mL HCl+100 mL alcohol)
was used as the etching agent. The porosity was analysed
on metallographic samples with an automatic image-
analysing system (AnalySIS-pro 3.0) built into the light
microscope (Microphot FXA, Nikon with 3CCD digital
camera Hitachi HV-C20A).
The local microchemical composition of the bearing
samples was analyzed with X-ray Energy Dispersion
Spectrometry (SEM FE JEOL JSM-6500F with EDXS
Inca Energy 400). The accelerating voltage of the
primary electron beam was 20 kV and an analysed
volume of approximately 1 μm3 was used for the selected
analyzing conditions. The EDXS method is not suitable
for a determination of the local carbon and oxygen
contents. It can serve as a qualitative estimate only. It
must also be warned here that incomplete removal of the
oil from a sample could seriously damage the electron
beam source of the SEM.
3 RESULTS AND DISCUSSION
A simple thermodynamic analysis makes it possible
to estimate the expected microstructures of the bearings.
In our case, computer tools for the prediction of the
phase equilibrium and the experimentally determined
phase diagrams were used8–10. The selected multi-com-
ponent (Cu-Sn-Pb) and binary phase diagrams (Cu-Sn,
Cu-Pb in Sn-Pb) show that in our system a substitutional
solid solution of Sn in Cu can be expected. The solubility
of Sn in Cu at room temperature is low. In the case of a
low cooling rate, therefore, besides the FCC_A1 solid
solution one can also expect Cu3Sn and the  phase.
B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
336 Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342
Figure 1: Snap-shot of bearing C (a), original magnification
6.7-times, and macro snap-shot of metallographic sample (b), Stereo
microscope Olympus SZ61
Slika 1: Makroskopski posnetek le`aja C: (a) originalna pove~ava
6,7-krat in makro posnetek metalografskega obrusa pri originalni
pove~avi 10-krat (b); Stereo mikroskop Olympus SZ61
Figures 2 a and b show the experimentally and theore-
tically calculated binary Cu-Sn phase diagrams. The
latter is almost completely in accordance with the
experimentally determined diagram. During the sintering
of the non-homogeneous mixtures, one can also expect
other hard and brittle intermetallic phases, such as
Cu6Sn5 and similar. Lead (Pb) is practically insoluble in
Cu and Sn and appears as a secondary phase. In the past,
it was added as an alloying element to improve the
functional properties of the bearings. Nowadays,
however, it is no longer used, for ecological and health
reasons.
A visual inspection showed that all the investigated
bearings have a golden metallic colour, typical for a
bronze bearing alloy. However, they differ slightly in
terms of the colour nuances, and also in terms of the
design (size and shape). The average sintered densities of
the oil-impregnated bearings determined with Mohr’s
balance are given in Table 1. Bearing A has the largest
density, but bearings B and C have similar values. It is
obvious that bearings B and C have a too low density,
when considering DIN 30910, Part 3.
Table 1: Density of sintered bearings
Tabela 1: Gostota sintranih le`ajev
Sample designation Density (g/cm3) Bearing weight (g)
Bearing A 7.03 1.49
Bearing B 6.67 0.93
Bearing C 6.64 0.88
The results of the average bulk chemical composition
of the investigated bearings are given in Table 2. Bearing
A has the highest content of Sn and the lowest Pb and Zn
contents. This could be a sign for the best sliding proper-
ties of this bearing. Bearings B and C have a much
higher content of Pb and Zn. This means that very
impure raw materials (powders) were used. The quantity
of the other possible elements (As, Ni, Sb, and Al) was
also checked with the ICP-AES analyzer, but the
quantity of these elements was below 0.01 %.
B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342 337
Figure 3: Microstructure of a polished sample in the middle of
bearing A, visible at two different magnifications; direction perpendi-
cular to ADC
Slika 3: Mikrostruktura poliranega vzorca le`aja A v sredini, vidna pri
dveh razli~nih pove~avah; obrus je izdelan v pre~ni smeri na smer
stiskanja
Figure 2: Equilibrium binary phase diagram of Cu-Sn with clearly
visible peritectic, eutectic, eutectoid and peritectoid reactions: a)
experimentally determined9 and b) obtained with a Thermo-Calc
calculation (enlarged region 0–50 % of Sn)10
Slika 2: Ravnote`ni binarni fazni diagram Cu-Sn s peritektsko,
evtektsko ter evtektoidno in peritektoidno reakcijo: a) eksperimen-
talno dolo~en9 in b) dobljen s Termo-Calc izra~unom10; pove~ano
obmo~je od ni~ do 50 % Sn9
Figures 3, 4 and 5 show the microstructures of
polished samples of bearings A, B and C, with the clearly
visible porosity, their shape and size distribution.
Bearings A have typical interconnected porosity with a
relatively appropriate pore size distribution. This enables
good filling of the pores with oil and good sliding pro-
perties of the bearing. Larger inclusions are not present in
the microstructure. Besides larger pores, mainly smaller
pores are present; all located at the particle boundaries.
Larger differences in the longitudinal and perpen-
dicular directions were not observed. Therefore, one can
conclude that bearings A are compacted and sintered
uniformly and appropriately. Bearings B have a much
more non-uniform porosity in comparison with bearings
A. They have a non-uniform density and porosity along
the compaction direction (see Figure 6). The bearings
are denser in thinner sections and in some regions the
interconnected porosity is interrupted. The bearings C
are very porous. A large number of very small pores are
present and the already polished microstructure reveals
some chemical inhomogeneity. The differences in the
polished microstructure and porosity can be clearly
recognized in Figures 7 a, b and c.
As was already mentioned, the porosity was also
analysed with an automatic image-analysing system. The
pore size distribution of the analysed bearings is similar
(Figures 8 and 9), but the porosity of bearing C was
significantly larger (approx.  = 20 %) than those of
bearings A and B (approx.  = 10 %). Figures 8 and 9
B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
338 Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342
Figure 4: Microstructure of a polished sample in the middle of
bearing B, visible at two different magnifications; direction perpen-
dicular to ADC
Slika 4: Mikrostruktura poliranega vzorca le`aja B, vidna pri dveh
razli~nih pove~avah, obrus je izdelan v pre~ni smeri na smer stiskanja
Table 2: Average bulk chemical composition of the investigated bearings
Tabela 2: Povpre~na kemijska sestava preiskovanih le`ajev
Sample
designation
Chemical composition (w/%)
Remarks
Cu Sn Pb Zn Oil content
Bearing A 88.2 10.9 0.03 < 0.005 0.62 Method of
determination
ICP-AES
Bearing B 89.8 9.5 0.07 0.09 0.58
Bearing C 89.2 9.8 0.08 0.06 0.72
Figure 6: Microstructure of polished bearing B along the ADC
direction, snapshot made with MIA technique
Slika 6: Mikrostruktura poliranega vzorca le`aja B vzdol` osi stiska-
nja, posnetek izdelan v MIA-tehniki
Figure 5: Microstructure of a polished sample in the middle of
bearing C at two different magnifications; direction perpendicular to
ADC
Slika 5: Mikrostruktura poliranega vzorca le`aja C v sredini, vidna pri
dveh razli~nih pove~avah; obrus je izdelan v pre~ni smeri na smer
stiskanja
show histograms of the pore number and size (area)
distribution in the selected size regions (1  0–50 μm2, 2
 50–100 μm2, 3  100–200 μm2, 4  200–300 μm2, 5
 300–400 μm2, 6–400–500 μm2, 7  500–1000 μm2, 8
 1000–2000 μm2, 9  2000–5000 μm2, 10 
5000–10000 μm2). The analysis was performed at
magnifications of 50, 100- and 200-times. However, only
the smallest magnification gives a statistically relevant
number of pores (more than 1500).
Figure 8 shows that in all three cases more than 75 %
of the pores are in the smallest size range (0−50 μm2),
but this represents only about 18 % to 25 % of the total
volume of porosity (Figure 9). It is evident that bearing
C has a larger number of the smallest pores, but a smaller
number of pores in the next size regions (between 50 and
300 μm2). But, again, bearings C have a relatively large
proportion of larger pores in the size regions between
400 μm2 and 2000 μm2. It should be noted here that this
statistical analysis of pore number and size distribution
can only serve for a qualitative assessment because it
was not performed on a statistically large enough
number of samples.
Figures 10, 11 and 12 show microstructures of
etched samples of bearings A, B and C. One can observe
a very nice monophase microstructure of bearing A
(Figure 10), without any secondary phases present. In
the microstructure, Cu twins typical for FCC structures
are visible.
The Cu-Sn powder used for bearings A was very
likely an atomized CuSn10 alloy powder with a carefully
selected particle size distribution. Smaller particles fill
the empty places between the larger ones. The sintering
process was performed correctly with a relatively fast
cooling in the final stage of the sintering. Oxide
inclusions are not present. This shows an appropriately
performed process with respect to the sintering
B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342 339
Figure 9: Histogram of the size of pore area (AF – area of pores)
inside the individual area size region.
Slika 9: Histogramski prikaz ploskovnega dele`a por po posameznih
velikostnih razredih (AF – povr{ina por)
Figure 7: A comparison of microstructures of polished samples
visible under LM: a) bearing A, b) bearing B and c) bearing C, parallel
direction.
Slika 7: Primerjava med mikrostrukturami poliranih vzorcev: a) le`aj
A, b) le`aj B in c) le`aj C, obrusi izdelani v vzdol`ni smeri, sredina
Figure 8: Histogram of pore-number distribution (NP – number of
pores) inside the individual area size region.
Slika 8: Histogramski prikaz dele`a {tevila por po posameznih
velikostnih razredih (NP – {tevilo por)
atmosphere. Sintering the Cu89-Sn11 alloy usually runs
at approx. 820 °C in a mixture of nitrogen (N2) and
hydrogen (H2), in the ratio 70 : 301.
B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
340 Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342
Figure 11: Microstructure of the alloy of an etched sample of bearing
B, visible at two different magnifications; perpendicular to ADC
Slika 11: Mikrostruktura zlitine jedkanega vzorca le`aja B, vidna pri
dveh razli~nih pove~avah; obrus je izdelan v pre~ni smeri
Figure 10: Microstructure of an etched sample of bearing A, visible at
two different magnifications; perpendicular to ADC
Slika 10: Mikrostruktura zlitine jedkanega vzorca le`aja A:, vidna pri
dveh razli~nih pove~avah, obrus je izdelan v pre~ni smeri
Figure 13: Comparison of the etched microstructures of alloys: a)
bearing A, b) bearing B and c) bearing C, perpendicular to the
compaction direction.
Slika 13: Primerjava med mikrostrukturami zlitin jedkanih vzorcev: a)
le`aj A, b) le`aj B in c) le`aj C, obrusi izdelani v pre~ni smeri, sredina.
Figure 12: Microstructure of the alloy of an etched sample of bearing
C, visible at two different magnifications; perpendicular to ADC
Slika 12: Mikrostruktura zlitine jedkanega vzorca le`aja C, vidna pri
dveh razli~nih pove~avah
The microstructure of the alloy of bearing B (Figure
11) is much more heterogeneous. Larger parts of the fine
pores and inclusions are present. Large fields of slightly
light pink and blue show the presence of secondary
phases and chemical inhomogeneity. This shows that a
powder mixture consisting of elemental Cu and Sn
particles was probably used.
The microstructure of bearing C (Figure 12) is also
very heterogeneous. One can observe differently co-
loured regions. Typical FCC twins are visible here, and
also regions with different chemical compositions. Inside
the larger grains one can also see small subgrains. The
differences in the etched microstructures of the present
bearings can be easily recognized in Figures 13 a to c.
In Table 3, the average Vickers hardness of the
investigated bearings is given. Bearings A have the
highest and the most uniform hardness values. The
hardness of the bearings B and C is much lower and
non-uniform. This is in accordance with the observations
of the microstructure under LM.
Table 3: Hardness measurements of the sintered bearings.
Tabela 3: Trdota izmerjena na metalografskih obrusih sintranih
le`ajev
Sample designation Measuring place HV1 /MPa
Bearing A
perpendicular 79.7
parallel 69.1
Bearing B
perpendicular 46.3
parallel 40.5
Bearing C
perpendicular 36.3
parallel 50.4
Scanning electron microscopy and mapping micro-
analyses (SEM/EDXS) were performed at different
locations of the investigated bearings, as shown in
Figures 14 and 15. In all cases Cu and Sn are mainly
present. However, in bearings B and C, Zn is also
B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342 341
Figure 15: SEI image of the analysed surface of bearing sample C with designated regions where microchemical analyses were made (a); and the
corresponding characteristic EDXS spectrum 2 with visible peaks for Cu, Sn, O (b)
Slika 15: SEI-posnetek analizirane povr{ine le`aja C z ozna~enimi mesti, kjer je bila opravljena ploskovna EDXS-mikrokemijska analiza (a) in
pripadajo~ karakteristi~en EDXS-spekter 2 z vidnimi vrhovi za Cu in O ter manj{im vrhom Sn (b)
Figure 14: SEI image of the analysed surface of bearing sample B with designated regions where the microchemical analyses were made (a); and
the corresponding characteristic EDXS spectrum 4 with visible peaks for Cu, Sn, O in Zn (b).
Slika 14: SEI-posnetek analizirane povr{ine le`aja B z ozna~enimi mesti, kjer je bila opravljena ploskovna mikrokemijska EDXS-analiza (a) in
pripadajo~ karakteristi~en EDXS-spekter 4 z vidnimi vrhovi za Cu, Sn, O in Zn (b)
present. The SEM/EDXS analyses of bearing A showed
its relatively uniform chemical composition. The local
concentration of Sn in the individual particles varied
between mass fractions 8.4 % to 11.4 %. It was
concluded that all the particles have practically the
selected bronze bearing-alloy composition. The assessed
content of oxygen is approx. 0.5 %.
The bearing B has a local content of Sn between mass
fractions 8.3 % and 12.7 %. At one location there was a
particle also containing Zn (3.6 %) and a small content
of Sn (1.9 %). Obviously, powder particles of brass are
also present in the powder mixtures used. Bearing C has
the lowest local content of Sn (between 6.1 % and
11.6 %). At two locations, particles of practically pure
Cu are found. This confirms that a mixture of elemental
powders was used. Particles containing Zn were not
found, although a bulk chemical analysis established its
presence. Obviously, also in this case powder particles of
brass are present in the powder mixtures. Bearings B and
C also have a higher oxygen content compared to the
bearings A. The assessed oxygen content of the bearing
B is between 0.6 % and 2.0 % and that of bearing C is
between 0.7 % and 1.2 %. This indicates that a poorer
sintering atmosphere was used, compared to the sintering
atmosphere used for bearings A.
4 CONCLUSIONS
This work presents an excellent example of how the
selected raw materials and process conditions can
drastically influence the final functional properties of a
sintered product. The results of our analyses and
investigations showed that all the investigated bearings
have a similar average chemical composition (bronze of
type Cu-10 % Sn). However, they significantly differ in
terms of the microstructure and consequently also in the
mechanical and functional properties. This is first of all
due to the different raw materials being used. It can be
supposed that for the production of bearing A a
chemically homogeneous Cu-Sn10 % atomized alloyed
powder is used, while for the production of bearings B
and C powder mixtures of elemental Cu and Sn are used.
In this case, less pure raw materials are also used with
some presence of Zn and Pb. In the case of bearing A
also much better conditions of compaction and the
sintering process were used. The density of bearing A is
more uniform and the interconnected porosity enables
optimal oil filling. These bearings exhibit no presence of
secondary intermetallic phases, oxides and other
inclusions. The average hardness of the bearings A is
much higher and uniform. All these facts explain why
these bearings have much better functional properties in
comparison with the bearings B and C.
5 REFERENCES
1 P/M-Encyclopedia; CD-ROM; European Powder Metallurgy Asso-
ciation, 2001
2 Sint - Prüfnormen (SPN); Sintermetalle; DIN 30 910, Teil 1&3,
WLB, Sintermetalle für Lager und Formteile mit Gleiteigenschaften,
Oktober 1990
3 Thümmler, R. Oberacker: An Introduction to Powder Metallurgy,
The Institute of Materials, The University Press, Cambridge, Lon-
don, UK, 1993
4 R. M. German: Powder Metallurgy Science, Metal Powder Industries
Federation (MPIF), Princeton, New Jersey, 2nd edition, 1994
5 E. Klar et al: Powder Metallurgy, Metals Hand Book, 9th Edition,
Volume 7, P/M ASM Handbook Committee, American Society For
Metals, Ohio, 1984
6 Sintered metal materials, excluding hardmetals – Permeable sintered
metal materials – Determination of density, oil content and open
porosity; International standard ISO
7 Sint - Prüfnormen (SPN); Prüfung der Sinterhärte; DIN 30 911, Teil
4, Oktober 1990
8 Database of Thermochemical and Physical Properties (TAPP), ES
Microware, Hamilton, Ohio, ZDA
9 E. A. Brandes, G. B. Brook: Smithells metals reference Book, 7th
Edition, Butterworth Heinemann, Oxford, 1992
10 ThermoCalc; web page: http://www.thermocalc.com/
B: [U[TAR[I^ et al.: THE INFLUENCE OF POWDER SELECTION AND PROCESS CONDITIONS ...
342 Materiali in tehnologije / Materials and technology 44 (2010) 6, 335–342
D. A. SKOBIR et al.: THE INFLUENCE OF THE MICROALLOYING ELEMENTS OF HSLA STEEL ...
THE INFLUENCE OF THE MICROALLOYING
ELEMENTS OF HSLA STEEL ON THE
MICROSTRUCTURE AND MECHANICAL PROPERTIES
VPLIV MIKROLEGIRNIH ELEMENTOV NA MIKROSTRUKTURO
IN MEHANSKE LASTNOSTI HSLA JEKLA
Danijela A. Skobir1, Matja` Godec1, Martin Balcar2, Monika Jenko1
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2@ÏAS, a. s., Strojírenská 6, 591 71 @d’ár nad Sázavou, Czech Republic
danijela.skobir@imt.si
Prejem rokopisa – received: 2010-08-30; sprejem za objavo – accepted for publication: 2010-09-21
In this work, hot-forged, high-strength, low-alloy (HSLA) steels based on the chemical composition 34CrNiMo6 were
investigated in order to determine the type of precipitates forming and their effect on the mechanical properties. The individual
steels were microalloyed with one of the following elements or a combination of elements: niobium, titanium or zirconium and
vanadium. The compositional changes in the microstructure and the various kinds of precipitates observed in the microalloyed
steels were examined using energy-dispersive spectroscopy (EDS). Optical and scanning electron microscopy studies revealed
that the addition of microalloying elements did not change very much the main microstructural features due to the fact that all
the microstructures consisted of tempered martensite and finely dispersed carbide precipitates along martensite laths. In all the
microstructures relatively large precipitates of the microalloyed elements were observed. This could be a possible reason for
observing no significant improvement in the mechanical properties.
Keywords: high-strength, low-alloy (HSLA) steels, microalloying elements, precipitates, mechanical properties, electron
microscopy
Vro~e kovano jeklo HSLA kemijske sestave 34CrNiMo6 smo raziskovali z namenom, da bi dolo~ili vrsto precipitatov, ki se
tvorijo, in njihov vpliv na mehanske lastnosti. Posamezna jekla so bila mikrolegirana z enim od naslednih elementov ali s
kombinacijo le-teh: z niobijem, s titanom ter cirkonijem in z vanadijem. Mikrostrukturne spremembe in razli~ne vrste
precipitatov smo {tudirali z metodo EDS. Preiskave s svetlobnim in vrsti~nim elektronskim mikroskopom so potrdile, da
dodatek mikrolegirnih elementov ni bistveno spremenil osnovne mikrostrukture, saj so bile vse mikrostrukture iz popu{~enega
martenzita z izlo~enimi karbidnimi precipitati po martenzitnih letvah. Pri vseh mikrostrukturah so bili opa`eni tudi precej veliki
precipitati mikrolegirnih elementov, ki so po vsej verjetnosti razlog za to, da se mehanske lastnosti niso bistveno izbolj{ale.
Klju~ne besede: visokotrdna malo legirana jekla (HSLA), mikrolegirni elementi, precipitati, mehanske lastnosti, elektronska
mikroskopija
1 INTRODUCTION
Microalloyed, high-strength, low-alloy (HSLA) steels
are technologically important structural materials.1,2
These steels contain small amounts of alloying elements,
such as titanium, niobium, vanadium or zirconium,
which enhance the strength through the formation of
stable carbides, nitrides or carbonitrides. Titanium also
readily forms stable nitrides at high temperatures and is
used to control the nitrogen content in the alloys. The
microalloying elements niobium and vanadium are added
to microalloyed cast steels primarily to provide grain
refinement and a response to aging. Niobium, often at
levels of less than 0.05 %, effectively prevents unde-
sirable grain growth and can also contribute to precipi-
tation strengthening. Vanadium, in particular at levels of
less than 0.1 %, forms strengthening carbonitride pre-
cipitates.3,4 Precipitation hardening by fine carbonitride
particles has been used for many years to increase the
strength of microalloyed steels. During the thermo-
mechanical processing, carbonitrides may nucleate in
austenite during forging, at the / interface during
transformation (interphase precipitation), or in super-
saturated ferrite during the final cooling.5
On adding such elements to steels with 0.008–0.03 %
C and up to w = 1.5 % Mn, it became possible to pro-
duce fine-grained material with yield strengths between
450 MPa and 550 MPa, and with ductile/brittle transition
temperatures as low as –70 °C.6
Nowadays, microalloyed cast steels have found many
applications in the manufacturing of industrial parts,
such as offshore platform nodes, centrifugal cast pipes,
machinery supports, ingot moulds and buckets, which all
used to be produced by expensive manufacturing pro-
cesses.7–9
2 EXPERIMENTAL WORK
Hot-forged HSLA steels based on the chemical
composition 34CrNiMo6 were investigated in order to
determine the type of precipitates forming and their
effect on the mechanical properties. Individual steels
were microalloyed with one of the following elements or
combination of elements: niobium (steel A), titanium
Materiali in tehnologije / Materials and technology 44 (2010) 6, 343–347 343
UDK 669.14.018.298: 620.1/.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)343(2010)
(steel B) or zirconium and vanadium (steel C). For com-
parison a non-microalloyed steel (steel D) was chosen.
The chemical compositions of the steels, as specified by
the manufacturer, are given in Table 1.
Metallographic samples were prepared using
standard polishing techniques and were then etched with
2 % Nital. The microstructure was examined in an opti-
cal microscope, a Nikon Microphot FXA. The carbide
precipitates were investigated using a field-emission
scanning electron microscope, a JEOL JSM 6500F,
equipped with an energy-dispersive spectroscopy (EDS)
analyzer (INCA X-SIGHT LN2 type detector, INCA
ENERGY 450 software). The accelerating voltage
during the EDS was 15 kV, the working distance 10 mm
and the probe current 0.05 nA.
To evaluate the mechanical properties tensile tests
were conducted on a 500-kN tensile machine, an Instron
1255.
3 RESULTS AND DISCUSSION
3.1 Microstructural characterization
The solubility data imply that in a microalloyed steel,
carbides and carbonitrides of Nb, Ti, Zr and V will
precipitate. While the primary effect of these fine
dispersions is to control the grain size, dispersion
strengthening will take place. The strengthening arising
from this cause will depend both on the particle size and
the interparticle spacing, which is determined by the
volume fraction of the precipitate. These parameters will
depend primarily on the type of compound that is
precipitating, and this is determined by the microalloying
content of the steel.6
In modern microalloyed steels there are at least three
strengthening mechanisms that contribute to the final
strength achieved. In the first mechanism, the precipi-
tation takes place in the austenite and further precipi-
tation occurs during the transformation to the ferrite. The
precipitation of niobium, titanium and vanadium
carbides has been shown to take place progressively as
the interphase boundaries move through the steel. This is
the second mechanism, called interphase precipitation.
As this precipitation is normally on an extremely fine
scale, occurring between 850 °C and 650 °C, it is likely
to be the major contribution to the dispersion strengthe-
ning. If the rate of cooling through the transformation is
high, leading to the formation of supersaturated plates of
ferrite, the carbides will tend to precipitate within the
grains, usually onto the dislocations, which are nume-
rous in this type of ferrite. This is the third strengthening
mechanism.6
Optical microscopy of the investigated steels revealed
that the addition of different microalloying elements did
not considerably change the main microstructural
features due to the fact that all the microstructures
consisted of tempered martensite and finely dispersed
carbide precipitates along the martensite laths. In all the
microstructures relatively large precipitates (up to 10
μm) of microalloyed elements were observed. After
etching, in the optical microscope these precipitates can
D. A. SKOBIR et al.: THE INFLUENCE OF THE MICROALLOYING ELEMENTS OF HSLA STEEL ...
344 Materiali in tehnologije / Materials and technology 44 (2010) 6, 343–347
Table 1: Chemical compositions of the investigated steels microalloyed with niobium (steel A), titanium (steel B), vanadium and zirconium (steel
C) and non-microalloyed (steel D) in mass fractions
Tabela 1: Kemijska sestava jekla, mikrolegiranega z niobijem (jeklo A), s titanom (jeklo B), z vanadijem in s cirkonijem (jeklo C) in nelegi-
ranega jekla (jeklo D) v masnih dele`ih (%)
Steel C Mn Si P S Cr Ni Cu Mo Al Nb
A 0.35 0.56 0.26 0.016 0.004 1.52 1.55 0.13 0.17 0.029 0.039
Steel C Mn Si P S Cr Ni Cu Mo Al Ti
B 0.32 0.55 0.27 0.022 0.002 1.58 1.56 0.15 0.18 0.020 0.027
Steel C Mn Si P S Cr Ni Cu Mo V Al Zr
C 0.34 0.54 0.27 0.017 0.002 1.58 1.55 0.21 0.23 0.030 0.031 0.030
Steel C Mn Si P S Cr Ni Mo Al
D 0.34 0.60 0.26 0.007 0.012 1.55 1.58 0.20 0.030
Figure 1: Optical micrographs of the steel microalloyed with niobium
(steel A) at different magnifications
Slika 1: Mikrostruktura jekla, mikrolegiranega z niobijem (jeklo A),
pri razli~nih pove~avah (svetlobni mikroskop)
be seen as typically coloured (brown-reddish for the
niobium and titanium precipitates and yellow for the
zirconium precipitates). Representative optical micro-
graphs of the microalloyed steels are presented in
Figures 1, 2 and 3.
In order to observe the microstructure more closely,
the precipitates of the microalloying elements were
investigated using scanning electron microscopy. In the
steel A, the precipitates are very similar to the degene-
rated eutectic Fe-Nb(C, N) known as "Chinese script"10,11
(Figure 4) that formed because of the improper solidi-
fication process of the cast ingots.
Figure 5 shows the large titanium nitride in the steel
B, bounded to a complex, non-metallic inclusion com-
posed of manganese sulphide and aluminium oxide.
A similar microstructure was also observed in the
steel C, where large zirconium carbonitrides were found
(Figure 6). EDS analyses of all three particles were
performed and the chemical composition is listed in
Table 2. Typical for the Nb, Ti and Zr carbonitrides is a
regular geometrical shape, which is seen as a triangular,
quadrilateral shape. Their shape is related to their crystal
structure and is influenced by growing along the ener-
getically most favourable crystal planes. Such particles
are not desirable in the material because they are very
hard and have sharp edges, which can cause a lowering
of the mechanical properties of the material.
The aim of microalloying the steel is to obtain very
fine Nb, Ti and Zr precipitates on the nanometre scale,
which improves the material properties with relatively
low costs. In our study it was found that the benefit of
microalloying was suppressed by a non-optimised
alloying process. We assumed that most of the micro-
alloying elements were extracted from the melt as large
particles instead of a finely dispersed precipitate.
Carbonitride precipitation has been the subject of
many investigations that have studied the different
aspects of the precipitates and their effects on the
mechanical properties.7 Based on these studies, the
effective carbonitrides in strengthening are those formed
in a fine form with an adequate distribution. As a result,
the increase observed in the strength can be attributed to
the precipitation of fine-scale carbonitrides in the form
of interphase and random precipitates.
D. A. SKOBIR et al.: THE INFLUENCE OF THE MICROALLOYING ELEMENTS OF HSLA STEEL ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 343–347 345
Figure 4: SE image and X-ray maps of niobium precipitate (steel A)
in the form of "Chinese script"
Slika 4: SE-posnetek in porazdelitev elementov niobijevega preci-
pitata v obliki kitajske pisave, dolo~ena z EDS
Figure 2: Optical micrographs of the steel microalloyed with titanium
(steel B) at different magnifications
Slika 2: Mikrostruktura jekla, mikrolegiranega s titanom (jeklo B), pri
razli~nih pove~avah (svetlobni mikroskop)
Figure 3: Optical micrographs of the steel microalloyed with vana-
dium and zirconium (steel C) at different magnifications
Slika 3: Mikrostruktura jekla, mikrolegiranega z vanadijem in cir-
konijem (jeklo C), pri razli~nih pove~avah (svetlobni mikroskop)
3.2 Mechanical properties
The results obtained from the tensile tests are
reported in Table 3. The properties reported in this table
are the average properties of two testing specimens.
According to these results, the presence of different
microalloying elements does not significantly change the
mechanical properties. It is possible that the benefit of
the fine precipitates is eliminated with the presence of
large Nb, Ti and Zr particles.
Table 3: Mechanical properties of the investigated steels
Tabela 3: Mehanske lastnosti preizku{anih jekel
Steel Re /MPa Rm /MPa A5/% Z/ %
A 611 805.5 18.6 53.1
B 637 806.5 18.7 64.6
C 627 808.5 19.2 65.2
D 619 786 21.1 62.2
4 CONCLUSIONS
The influence of the microalloying elements in
HSLA steels on the microstructure and mechanical
properties was investigated. It was found that Nb, Ti and
Zr precipitate not only as fine particles but also as large
particles with sharp edges. Based on the mechanical
properties and the SEM/EDS analysis it is supposed that
the benefit of microalloying elements is negated by these
large particles and the results are almost the same
mechanical properties as in the case of the non-micro-
alloyed steel (steel D).
5 REFERENCES
1 T. Gladman, The physical metallurgy of microalloyed steels. Lon-
don: Institute of Materials, (1997)
D. A. SKOBIR et al.: THE INFLUENCE OF THE MICROALLOYING ELEMENTS OF HSLA STEEL ...
346 Materiali in tehnologije / Materials and technology 44 (2010) 6, 343–347
Figure 5: SE image and X-ray maps of the titanium precipitate in the steel B
Slika 5: SE-posnetek in porazdelitev elementov titanovega precipitata v jeklu B
Figure 6: SE image of the zirconium precipitates in steel C
Slika 6: SE-posnetek cirkonijevih precipitatov v jeklu C
Table 2: Chemical composition for the analysed spectra in mass fract-
ions
Tabela 2: Kemijska sestava analiziranih spektov v masnih dele`ih (%)
C N Ti Cr Mn Fe Ni Zr
Spectrum 1 10.80 11.50 0.31 0.61 6.75 70.03
Spectrum 2 15.32 9.22 0.40 7.20 67.86
Spectrum 3 10.85 6.53 0.81 0.83 0.75 31.49 0.47 48.26
2 Z. K. Liu, Scripta Mater. 50 (2004), 601–606
3 S. N. Prasad, D. S. Sarma, Mater. Sci. Eng., A 399 (2005), 161–172
4 H. J. Kestenbach, S. S. Campos,d E. V. Morales, Mater. Sci.
Technol., 22 (2006), 615–626
5 M. D. C. Sobral, P. R. Mei, H. J. Kestenbach, Mater. Sci. Eng. A,
367, (2004), 317–321
6 H. K. D. H. Bhadeshia, R. Honeycombe, Steels Microstructure and
Properties, Third Edition, Elsevier (2006)
7 H. Najafi, J. Rassizadehghani, S. Asgari, Mater. Sci. Eng. A, 486,
(2008), 1–7
8 R. C. Voigt, M. Blair, J. Rassizadehghani, Proceedings of Interna-
tional Conference on 1990 Pressure Vessels and Piping, vol. 201,
ASME, Nashville, TN, USA, June 1990, 147–154
9 D. L. Albright, S. Bechet, K. Rohrig, Proceedings of International
Conference on Technology and applications of HSLA steels, ASM,
Philadelphia, PA, USA, 1984, 1137–1153
10 F. Haddad, S. E. Amara, R. Kesri, S. Hamar-Thibault, Journal de
Physique IV, 122 (2004), 35–39
11 J. Berneti~, B. Brada{kja, G. Kosec, B. Kosec, Mater. Technol. 42
(2008), 291–294
D. A. SKOBIR et al.: THE INFLUENCE OF THE MICROALLOYING ELEMENTS OF HSLA STEEL ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 343–347 347

G. KOSEC et al.: EFFECT OF 3 S REHEATING ON CHARPY NOTCH TOUGHNESS AND HARDNESS ...
EFFECT OF 3 S REHEATING ON CHARPY NOTCH
TOUGHNESS AND HARDNESS OF MARTENSITE AND
LOWER BAINITE
VPLIV 3 S POGREVANJA NA CHARPYJEVO ZAREZNO @ILAVOST
IN TRDOTO MARTENZITA IN SPODNJEGA BAINITA
Gorazd Kosec1, Franc Vodopivec2, Anton Smolej3, Jelena Vojvodi~-Tuma2,
Monika Jenko2
1SIJ Acroni, Jesenice
2Institute of Metals and Technology, Lepi pot 11, Ljubljana, Slovenia
3Faculty of Natural Sciences and Technology, A{ker~eva 6, Ljubljana, Slovenia
franc.vodopivec@imt.si
Prejem rokopisa – received: 2010-02-10; sprejem za objavo – accepted for publication: 2003-03-10
The 0.1C 0.032Nb fine grained steel was heat treated to fine and coarse grained lower bainite and martensite. Half of specimens
was reheated for 3 s at 750 °C with current conduction. Charpy notch tests were performed in temperatrure range from –200 °C
to 60 °C, room temperature hardness determined and microstructure and fracture surface examined with SEM. Notch toughness
and transition temperature are similar for lower bainite and delivered steel and much lower for martensite. After reheating, noth
toughness was decreased for about ten times for lower bainite and the cleavage temperature increased for about 80 °C. Both
changes were much lower for martensite and for as delivered steel. The effect of reheating on hardness was different for
different microstructure and much smaller as for notch toughness.
Key words: microalloyed structural steel, lower bainite, martensite, reheating, notch toughness, cleavage temperature, hardness
V drobnozrnatem jeklu z 0.1 C in 0.32 Nb je bila s toplotno obdelavo ustvarjena mikrostruktura iz martenzita in spodnjega
bainita. Polovica preizku{ancev je bila nato zmova segreta 3 s pri 750 °C z elektri~nim tokom. Charpy preizkusi so bili izvr{eni
v razponu temperature od –200 °C do +60 °C, trdota izmerjena pri sobni temperaturi, mikrostruktura in prelomne povr{ine pa
pregledane v SEM.. Zarezna `ilavost in prehodna temperatura sta podobni za dobavljeno jeklo in spodnji bajnit in ni`ji za
martenzit. Po ponovnem segretju je bila zarezna `ilavost zmanj{ana za okoli 10 krat, temperatura cepilnega preloma pa
pove~ana za okoli 80 °C za spodnji bainit. Obe spremembi sta bili mnogo manj{i pri martemnzitu. Vpliv ponovnega segravanja
je bil manj{i pri trdoti in razli~en pri razli~ni mikrostrukturi.
Klju~ne besede: mikrolegirano jeklo, spodnji bainit, martenzit, ponovno segrevanje, zarezna `ilavost, temperatura cepilnega
preloma, trdota
1 INTRODUCTION AND AIM OF THE WORK
In the heat affected zone (HAZ) od welds of
structural steels local brittle zones (LBZ) could form and
decrease the local toughness 1–13. Field experience and
laboratory tests suggested that for a 490 MPa fine
grained microalloyed steel with Charpy notch toughness
of about 250 J, lower bainite could be more propensive
to the formation of LBZ than martensite 14 and that the
HAZ propensity to embrittlement was for this steel
greater than for a conventional steel with the yield stress
of 350 MPa 15. Martensite and lower banite were
assumed to be more sensistive to the formation of LBZ
than other constituents of the HAZ microstructure an in
this work this assumption was verified and the effect of
grain sizeon Charpy toughness and transition tempera-
ture, rsp. cleavage temperature, and hardness was
investigated.
2 EXPERIMENTAL WORK
The investigation was carried ou with the high
strength low alloyed (HSLA) structural steel with
0.1C-0.5Mn-0.025Al-0.27Mo-0.032Nb-0.15Ni, the ini-
tial microstructure of fine grained ferrite and cementite
particles and the yield stress of 490 MPa. Austenite grain
size similar to that found in HAZ in contact with the
deposed metal was obtained with 3–4 s of heating of
Materiali in tehnologije / Materials and technology 44 (2010) 6, 349–355 349
UDK 669.14.018.62:621.791.05 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)349(2010)
Figure 1: Reheating cycle for specimens
Slika 1: Diagram ponovnega segrevanja preizku{ancev
single Charpy specimens at 1250 °C. These specimens
and specimens annealed for 20 min. at 920 °C were
quenched half in water at 70 °C and half in lead bath at
400 °C and martensite and lower bainite with different
grain size obtained. Half of specimens was than reheated
for 5 s at 750 °C with direct conduction heating and air
cooled according to Figure 1 in the heating part of a hot
deformation simulator. On all specimens the Charpy
notch was cut out after the heat treatment. The effect of
testing temperature in the range of -200 °C to 60 °C on
fracturing energy was than determined and the results for
as delivered steel used for comparison. The micro-
structure and of the brittle fracture surface were
investigated with optical and SE microscopy.
3 MICROSTRUCTURE
The constituents of microstructure formed at cooling
from 1250 °C and 920 °C are termed as primary and as
secondary the constituents formed at cooling from 750
°C. The fine grained microstructure of the as delivered
steel consisted mostly of acicular ferrite with small
stringers of fine cementite particles (Figure 2). After
reheating at 750 °C, small martensite inserts at triple
points and rare secondary martensite platelets were
found in the interior of ferrite grains. After quenching
from 920 °C in water a microstructure of martensite
platelets in ferrite matrix was obtained (Figure 3) that
changed after reheating to intergranular inserts and rare
intragranular platelets of secondary martensite (Figure
4). The cooling in lead bath from 920 °C produced a
microstruture with ferrite laths and stringers of cementite
particles. After reheating, the microstructure was similar
than for reheated martensite, however, it had more
frequent intragranular platelets and grain boundary
inserts of secondary martensite. After quenching from
1250 °C in water the microstructure consisted of the
same constituents as after cooling from 920 °C but with
a greater share of martensite, while, the grain size
coarser for 3 to 4 ASTM grades. After reheating, ferrite
platelets had frequently boundaries marked with
stringers of fine carbide particles and grain boundary
inserts of secondary martensite. After lead cooling from
1250 °C the microstructure consisted of platelets of
ferrite with boundary marked with cementite particles
(Figure 5). The lower bainite was not investigated to a
sufficient detail to conclude on the mechanism of
transformation, displacive or reconstructive 16,17,18. After
reheating, it changed to a microstructure of platelets of
secondary martensite and ferrite inside and inserts of
secondary martensite at boundaries of coarce grains
(Figure 6).
The short reheating at 750 °C and air cooling
produced the following changes of initial microstructures
of ferrite+bainite and ferrite+martensite:
– stringers of cementite particles in bainite were
dissolved producing platelets of secondary austenite
rich in carbon that transformed at cooling to platelets
of secondary martensite in the interior of ferrite
grains;
– primary martensite platelets were partially decom-
posed to ferrite and carbide precipitates;
– triple grain points and grain boundary inserts of
secondary austenite were formed and at cooling
transformed to secondary martensite. The size of
these inserts suggests a fast transport of carbon
towards the nucleation points of secondary austenite
at triple grains points It is possible that the
nucleation and growth of inserts were enhanced by
G. KOSEC et al.: EFFECT OF 3 S REHEATING ON CHARPY NOTCH TOUGHNESS AND HARDNESS ...
350 Materiali in tehnologije / Materials and technology 44 (2010) 6, 349–355
Figure 4: Microstructure after reheating
Slika 4: Mikrostruktura po ponovnem segrevanju
Figure 3: Microstructure after water quenching from 920 °C
Slika 3: Mikrostruktura po kaljenju v vodi z 920 °C
Figure 2: Microstructure of as delivered steel
Slika 2: Mikrostruktura uporabljenega jekla
segregation of carbon atoms at primary austenite
grains.
4 CHARPY TOUGHNESS AND TRANSITION
TEMPERATURE
In Figures 7 to 11 Charpy toughness is shown in
dependence of testing temperature. The upper shelf
toughness is high for the as delivered steel and for steel
with the microstructure of lower bainite with small effect
of grain size, as after cooling from 1250 °C the upper
shelf value is of about 220 J and after cooling from 920
°C it is of about 250 J. The upper shelf toughness tem-
perature was for martensite before and after reheating
above 60 °C. The lower shelf notch toughness is very
similar for all specimens, as found ealier 19 for a number
of structural steels. In the frame of accuracy of tests it
was not possible to determine an eventual effect of
microstructure on fracturing energy in clevage range.
With ecception of the as delivered steel, upper shelf
toughness, cleavage and transition temperature (half
upper shelf toughness temperature for high upper shelf
energy), were strongly affected by the change of micro-
structure after reheating and the changes were particu-
larly great for lower bainite. In all cases, the upper shelf
toughness was lower and the transition temperature was
higher for the reheated steel.
G. KOSEC et al.: EFFECT OF 3 S REHEATING ON CHARPY NOTCH TOUGHNESS AND HARDNESS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 349–355 351
Figure 9: Dependence Charpy toughness versus testing temperature
after cooling from 920 °C in lead bath at 400 °C and after reheating at
750 °C
Slika 9: Odvisnost med Charpyjevo `ilavostjo in temperaturo preiz-
ku{anja za jeklo ohlajeno z 920 °C v svincu in po ponovnem
segrevanju pri 750 °C
Figure 6: Microstructure after reheat
Slika 6: Mikrostruktura po ponovnem segrevanju
Figure 5: Microstructure after lead bath quenching from 1250 °C
Slika 5: Mikrostruktura po kaljenju v vodi s 1250 °C
Figure 8: Dependence Charpy toughness versus testing temperature
after hot water quenching from 920 °C and after reheating at 750 °C
Slika 8: Odvisnost med Charpyjevo `ilavostjo in temperaturo preiz-
ku{anja za jeklo kaljeno z 920 °C v vro~i vodi in po ponovnem
segrevanju pri 750 °C
Figure 7: Dependence Charpy toughness versus testing temperature
for the steel with the as delivered steel and after reheating at 750 °C
Slika 7: Odvisnost med Charpyjevo `ilavostjo in temperaturo preiz-
ku{anja za dobavljeno jeklo in po ponovnem segrevanju pri 750 °C
After reheating the as delivered steel, the transition
temperature increased from –67 °C to –3 °C (Figure 7)
and the upper shelf toughness lovered from 240 J to 200
J. After quenching in water from 920 °C, notch
toughness at 0 °C was diminished to one half and after
reheating to one fourth of that for the as delivered steel
(Figure 8). The upper shelf range of the water quenched
steel was not achieved at 60 °C. With comparison to
water cooling from 920 °C, after higher austenitising
temperature notch toughness was affected stronger, the 0
°C energy was lower for about 2.5 to 3 times after
quenching and reheating, while the clevage temperature
was virtually not affected (Figure 9).
The Charpy toughness of about 250 J and the
transition temperature of –80 °C were obtained for steel
cooled from 920 °C in lead bath (Figure 10). After
reheating, the Charpy level was diminiseh the must of all
tested cases, at 0 °C by aproximately ten times, from 250
J to about 25 J. The cleavage temperature was increased
from about –100 °C to 0 °C and at 60 °C the Charpy
energy was still about four times smaller than that for the
investigated steel. The upper shelf energy was lower for
about 35 J after cooling from the higher temperature and
the clevage and transition temperature increased for
about 40 °C (Figure 11). The effect of reheating was
even stronger that after cooling from 920 °C and 0 °C
energy lowered for about 20 times and the the cleavage
temperature higher for about 80 °C (Figure 13).
These results indicate that independently on grain
size lower bainite is much nore prone to embrittlement
after short reheat at 750 °C than the steel with the
microstructure of small grained ferrite and pearlite as
well of fine and coarse grained martensite.
5 HARDNESS
In Table 1 hardness is shown for specimens with
different microstructure. The lowest hardness of the as
delivered steel increased significantly after reheating.
After water quenching from 920 °C the hardness was
increased greatly and it was lower after reheating. After
quenching from 920 °C in lead bath a relatively low
hardness was obtained, which was even lower after
reheating. After water quenching from 1250 °C the
greatest hardness was obtained that diminished
significantly after reheating, still, remaining high. After
quenching in lead bath from 1250 °C the hardness was
increased moderately in comparison to the as delivered
steel and it was higher after reheating.
The hardness level corresponds to the microstructure
after cooling and changes of hardness after reheating
agree with changes of microstructure. The effect of
reheating is for hardness in most cases lower than for
notch toughness. The changes of both properties are not
allways correlated, since in some cases by lower
hardness notch toughness is lower, also.
Table 1: Hardness of steel after different thermal treatmen
Tabela 1: Trdota jekla po razli~ni toplotnim obdelavi
Thermal treatment Hardness HV 5
As delivered 205
As delivered + 750 °C 248
920 °C  water 282
920 °C  water + 750 °C 244
920 °C  lead bath 222
920 °C  lead bath + 750 °C 214
1250 °C  water 383
1250 °C  water + 750 °C 320
1250 °C  lead bath 257
1250 °C  lead bath + 750 °C 298
6 FRACTURE SURFACE
Three fracturing mechanisms were identified for the
three levels of consumed energy. For high fracturing
G. KOSEC et al.: EFFECT OF 3 S REHEATING ON CHARPY NOTCH TOUGHNESS AND HARDNESS ...
352 Materiali in tehnologije / Materials and technology 44 (2010) 6, 349–355
Figure 11: Dependence Charpy toughness versus testing temperature
after cooling from 1250 °C in lead bath at 400 °C and after reheating
at 750 °C
Slika 11: Odvisnost med Charpyjevo `ilavostjo in temperaturo preiz-
ku{anja za jeklo ohlajeno z 1250 °C v svincu in po ponovnem
segrevanju pri 750 °C
Figure 10: Dependence Charpy toughness versus testing temperature
after hot water quenching from 1250 °C and after reheating at 750 °C
Slika 10: Odvisnost med Charpyjevo `ilavostjo in temperaturo preiz-
ku{anja za jeklo kaljeno z 1250 °C v vro~i vodi in po ponovnem
segrevanju pri 750 °C
energy, the fracture surface was of uneven profile and
consisted of normal and areas of shear decohesion
(Figure 12 and 13) 19 with dimples of different shape
and size. A specific fracture surface was observed on
specimens fractured with low energy consumption in the
range of temperature of growth of fracturing energy
above the cleavage threshold. It consisted of a mixture of
brittle and ductile morphology with prevalence of brittle
fracturing for low consumed energy (Figure 14) and an
increased share of of ductile fracturing for greater
consumed energy. In flat part the boundary of cleavage
propagation area was not clear and it is assumed that the
transition from brittle to ductile propagation occurred
with plane slip 20.
In lower shelf range on the fracture of the as received
steels, the shape and size of brittle facets was dependent
on the size of ferrite or austenite grains (Figure 15).
Generally, on the fracture surface of specimens cooled to
lower bainite, a greater number of rivers – ligaments of
microcracks propagating at different level of the same
cleavage plane was observed (Figure 16). After
quenching for 1250 °C the fracture constituents were
coarser because of the greater austenite grain size and for
lower bainite again more rivers and a more fragmented
fracture facets were found. In presence of martensite and
ferrite platelets, the crak propagates with coalescence of
microcracks in grains wit different lattice orientation and
microcracks in parallel lattice planes index join in rivers
that did not mark the croissing of fracture from ferrite to
martensite area. Also, no particular fracture details were
found that could be assotiated with the presence of grain
boundary inserts of secondary martensite in reheated
steels. No difference in clevage morfphology was found
between specimens with low and increaseed clevage
G. KOSEC et al.: EFFECT OF 3 S REHEATING ON CHARPY NOTCH TOUGHNESS AND HARDNESS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 349–355 353
Figure 16: Steel quenched from 1250 °C in lead and fractured at –60
°C
Slika 16: Jeklo ohlajeno s 1250 °C v svincu in prelomljeno pri –60 °C
Figure 14: Steel quenched from 920 °C in water and fractured at 0 °C
Slika 14: Jeklo kaljeno z 920 °C v vodi in prelomljeno pri 0 °C
Figure 13: As delivered steel, fracture surface at 22 °C. Shear decohe-
sion
Slika 13: Dobavljeno jeklo, prelomna povr{ina pri 22 °C. Cepilna
dekohezija
Figure 12: As delivered steel, fracture surface at 22 °C
Slika 12: Dobavljeno jeklo, prelomna povr{ina pri 22 °C
Figure 15: Steel quenched in water from 920 °C and fractured at –60
°C
Slika 15: Jeklo kaljeno v vodi z 920 °C in prelomljeno pri –60 °C
temperature. According to 21 the number of rivers is
greater by crack propagation in (110) than in (100)
ferrite lattice planes. EBSD examination of clevage
facets has shown that for the investigated microstructures
clevage occurred only in the (100) lattice plane
(Figure17) 15.
7 DISCUSSION
Most of the fracturing energy in ductile range is
consumed for the plastic deformation before the crack is
started at the notch tip and it is dissipated as adiabatic
heat 22. For this reason, the fracturing temperature at the
crack tip is higher than the nominal test temperature and
both temperatures are equal only below the cleavage
threshold temperature,while, in transition and ductile
fracturing range the difference increases with greater
plastic deformation in the limited volume of steel
involved in the deformation and fracturing events 22. In
this discussion it is assumed that the fracturing tempe-
rature is equal to the nominal temperature, which is
given as abscissa in Figure 7 to 11 and in the captions of
microfractographies.
In lower shelf range, where the fracturing occurs after
elastic deflection with a consumption of 5 J to 7 J 19, the
eventual effect of microstructure was below the level of
sensitivity of performed Charpy tests. For the same steel
the fracturing energy is different for different grain size
and the transition temperature is much lower for lower
bainite than for martensite, while after reheating, the
fracturing energy is decreased and the transition
temperature increased much more for lower bainite. The
higher sensistivity of lower bainite to reheating reflects
the effect of changes of microstructure. As for martensite
and lower bainite inserts of secondary martensite are
found after reheating, it is clear that the extension of
cleavage range is due to intragranular changes, speci-
fically the presence of secondary martensite platelets
formed with dissolution of cementite particles in secon-
dary austenite. The nature of changes in lower bainite
because of short reheating and changes in martensite at
short reheating that lower notch toughness are inve-
stigated. The so far obtained experimental findings
indicate that at reheating temperature no direct transfor-
mation of martensite to austenite took place and that rate
of dissolution of carbon in ferrite resp. scondary
austenite is faster that the precipitation of iron carbide
from the solid solution in martensite.
For some microstructures the changes of notch
toughness and hardness are as expected: by increased
hardness notch toughness is lower and opposite. Two
cases deviate significantly for the general rule: hardness
and notch toughness are lower for martensite after
reheating, while by very small difference in hardness
notch toughness is much lower for lower bainite. This
suggests that in weld heat affected zone lower hardness
is not always related to higher notch toughness.
8 CONCLUSIONS
On the base of experimental findings and their ana-
lysis the following conclusions are proposed:
– independently on grain size, Charpy notch toughness
is much higher and the transition temperature is
much lower in case of transformation of coarse and
fine grained HAZ austenite to bainite than to marten-
site;
– inpendently on grain size, after short reheating at
750 °C, Charpy notch toughness is greatly dimi-
nished for lower bainite, while, it is only slightly
diminished for martensite;
– particularly deleterious for notch toughness and
transition temperature is the formation of secondary
martensite from secondary austenite formed with
dissolution of cementite particles in interior of grains
at reheating;
– depending on microstructure, by similar hardness
different notch toughness and transition temperature
can be obtained for different microstructure;
– although beneficial in terms of notch toughness and
transition temperature, lower bainite in heat affected
G. KOSEC et al.: EFFECT OF 3 S REHEATING ON CHARPY NOTCH TOUGHNESS AND HARDNESS ...
354 Materiali in tehnologije / Materials and technology 44 (2010) 6, 349–355
Figure 17: (a) SE micrography of the fracture surface of the steel
quenched from 1250 °C in water at 70 °C and reheated at 750 °C; (b)
EBSD analysis and indexing of a cleavage facet 23
Slika 17: (a) SE posnetek prelomne povr{ine jekla, ki je bilo kaljeno s
1250 °C v vodi pri 70 °C in ponovno segreto pri 750 °C; (b) EBSD
analiza in indeksi cepilne ravnine
zone of welds increases more the sensibility to
formation of LBZ than martensite.
9 REFERENCES
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Intern. Conf. The Metallurgy, Welding and Qualification of Micro-
alloyed (HSLA) Steels Weldments, American Welding Society, 1990,
351–382
2 T. M. Sconover. Evaluation of Local Brittle Zones in HSLA-80
Weldments: Proceed. of The Intern. Conf. The Metallurgy, Welding
and Qualification of Microalloyed (HSLA) Steels Weldments,
American Welding Society, 1990, 276–305
3 S. Aihara, K. Okamoto: Influence of Local Brittle Zones on HAZ
Toughness of TMCP Steels: Proceed. of The Intern. Conf. The
Metallurgy, Welding and Qualification of Microalloyed (HSLA)
Steels Weldments, American Welding Society, 1990, 401–425
4 Q. Liu, X Diao, Y. W. Mai: Engineering Fracture Mechanics 49
(1994), 741–750
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6 T. Moltubak, C. Thaulow, Z. L. Zhang: Engineering Fracture Mecha-
nics 62 (1999), 445–462
7 J. Vojvodi~-Tuma, A. Sedmak: Analysis of the unstable fracture
behaviour of a high strength low alloy steel weldment: Engineering
Fracture Mechanics 71 (2004), 1435–1451
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(2007), 2395–2419
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(1994), 741–749
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13 J. Il Jang, B. W. Lee, J. B. Ju, D. Kwon, W. S. Kim: Engineerimng
Fracture Mechanics 70 (2003), 1245–1257
14 F. Vodopivec, B. Arzen{ek, D. Kmeti~, J. Vojvodi~-Tuma: Mater.
Tehnol. 37 (2003), 317–326
15 H. K. D. K Bhadesha, D. V. Edmonds: Acta Metall. 28 (1980),
1265–1273
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17 H. Matsuda, H. K. D. H. Bhadesha: Proc. R. Soc. Lond. A (2004),
460, 1707–1722
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steels with different microstructure; Proceedings of the Intern. Conf.
Mechanical Properties of Advanced Engineering Materials, ed. M.
Tokuda, B. Xu, Mie Un. Japan and Tsingua Un. China, 2001,
187–200
19 G. Kosec: Krhki prelom v coni toplotnega vpliva zvarov jekla
Niomol 490K (Brittle fracturing in heat affected zone of steel
Niomol 490K; dr. thesis (in Slovene), University of Ljubljana, 2007
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Tuma: Mat. Sci. Techn. 18 (2002), 68–72
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47 (2008), 173–179
The authors are indebted to the Ministry of High
Education, Science and Technology of Slovenia, Ljub-
ljana and the company SIJ Acroni, d. o. o. Jesenice for
the support of this investigation.
G. KOSEC et al.: EFFECT OF 3 S REHEATING ON CHARPY NOTCH TOUGHNESS AND HARDNESS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 349–355 355

S. TADI] et al.: DEFORMATION MECHANISMS IN Ti3Al-Nb ALLOY AT ELEVATED TEMPERATURES
DEFORMATION MECHANISMS IN Ti3Al-Nb ALLOY AT
ELEVATED TEMPERATURES
MEHANIZEM DEFORMACIJE V ZLITINI Ti3Al-Nb PRI
POVI[ANIH TEMPERATURAH
Srdjan Tadi}1, Radica Proki}-Cvetkovi}1, Igor Bala}1, Radmila
Heinemann-Jan~i}2, Katarina Boji}1, Aleksandar Sedmak1
1Innovation Centre, Faculty of Mechanical Engineering, University of Belgrade, Serbia
2 Faculty of Technology and Metallurgy, University of Belgrade, Serbia
srdjantadic26@gmail.com
Prejem rokopisa – received: 2010-09-10; sprejem za objavo – accepted for publication: 2010-11-20
Ti3Al-11Nb-1V (x/%) was made in a vacuum arc furnace. Hot-rolling was conducted at 1050 °C in a two-phase field of the
alloy. Strain-rate jump tests were performed at T = 800–1100 °C and strain rates of 10–6 to 10–3 s–1. Creep stress exponent (n)
and activation energy for deformation (Q) were used to identify deformation mechanisms. Two sequential mechanisms are
revealed: (i) viscous glide, typical of solid solution alloys and, (ii) power law breakdown at lower temperatures. Activation
energy for deformation was found to be strongly stress-dependant.
Keywords: Titanium intermetallics, deformation, viscous glide, power law breakdown
Zlitina Ti3Al-11Nb-1V (x/%) je bila izdelana v vakuumski oblo~ni pe~i. Vro~e valjanje je bilo opravljeno pri 1050 °C v
dvofaznem obmo~ju zlitine. Sunkoviti preizkusi hitrosti deformacije so bili narejeni pri T = 800–1100 °C in pri hitrosti
deformacije 10–6 to 10–3 s–1. Eksponent napetosti lezenja (n) in aktivacijska energija deformacije (Q) sta bila uporabljena za
identifikacijo mehanizma deformacije. Odkrita sta bila dva sekven~na mehanizma: (i) viskozno drsenje, ki je zna~ilno za zlitine
iz trdnih raztopin in (ii) prelom poten~nega zakona pri ni`ji temperaturi. Ugotovljeno je bilo, da je aktivacijska energija
deformacije zelo odvisna od napetosti.
Klju~ne besede: titanova intermetalna zlitina, deformacija, viskozno drsenje, poten~ni zakon, prelom zakona
1 INTRODUCTION
High-temperature mechanical properties and low
density make the titanium aluminides attractive candi-
dates for engine and airframe applications. Considerable
efforts in research have been devoted to this class of
materials, particularly to Ti3Al-Nb alloy. This alloy
consists typically of the ordered Ti3Al 2 phase (DO19
structure, similar to h. c. p) and the b. c. c.  solid solu-
tion phase 1.
The main handicap of Ti3Al-Nb alloys at room
temperatures is attributed to limited ductility. It is
ascribed due to the inadequate number of crystallo-
graphic slip systems 2. However, at a higher temperature
the alloy exhibits good ductility through thermal
activated softening 3. The high-temperature plasticity,
thermomechanical processing and deformation mecha-
nisms attracted the attention of a number of authors 4.
Three distinct high-temperature regimes of the thermo-
mechanical processing were reported 5: warm working
(below 950 °C), hot working in the two phase field
(980–1040 °C) and a hot working regime above the
-transus. In a more complex Ti-25Al-10Nb-3V-1Mo
alloy, dynamic recovery and dynamic recrystallization
were identified in the two-phase (2+) field 6. It was
suggested that the activation barrier for DRX in the
(2+) field depends on the cross-slip in the 2 phase 7.
Creep 8 and superplastic deformation 9 has also been well
examined, although a lot of uncertainties remained. The
identification of the rate controlling deformation
mechanisms is still not fully clear in two phase alloys.
Particularly confusing is the large scatter of activation
energies 10,11 associated with deformation processes.
Creep, transition to low-temperature creep and power
law breakdown behaviour is the emphasis of this work.
Deformation mechanisms are identified in terms of the
creep stress exponent and activation energy for
deformation.
2 EXPERIMENTAL
The alloy was prepared from a Ti-6Al-4V (with an
extra low interstitials quality), high-purity aluminium
and niobium flakes. It was smelted in a vacuum arc
furnace under a low-pressure 'argon blanket'. Remelting
was repeated several times in order to obtain a
homogeneous, pore-free rod 20 mm in diameter. The
nominal chemical composition of the alloy was as
follows (by x/%): 24 % Al, 11 % Nb and 1 % V, Ti – the
rest. After cleaning, the rod was hot-rolled on a labo-
ratory two-high mill with 	-calibrated rolls, down to a
diameter of 6 mm. Hot-rolling was conducted at 1050 °C
in the two-phase field, well below the -transus tempe-
rature. The rolling direction was swapped between two
successive passes to eliminate the temperature gradient.
Materiali in tehnologije / Materials and technology 44 (2010) 6, 357–361 357
UDK 669.295:621.77:539.37/.38 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)357(2010)
Compression test specimens with 	 = 6 mm and 12
mm in height were machined from the previously
chemically cleaned rod. The specimens were coated with
fine molybdenum-sulfide powder to minimise barrelling.
A series of strain-rate jump tests 12 were conducted on an
"Instron" testing machine in a water-cooled vacuum
chamber at temperatures ranging from 800 °C to 1100
°C. The temperature of the chamber was maintained
within 
5 °C. The true strain rates were calculated 13 to
be in the range from 10–6 to 10–3 s–1. The tests were con-
ducted down to 50 % of reduction of the specimen
height. True stress was normalized with shear modulus 
= 5.1 × 103 – 9T 14.
Metallographic specimens were prepared by wet
grinding and diamond-paste polishing down to 1/4 μm.
Keller’s etchant was used to reveal the microstructure
under the optical microscope.
3 RESULTS
The microstructure of the hot-rolled bar in a view
normal to the rolling direction is shown in Figure 1. The
microstructure consists of a homogeneously distributed
2 phase in the  matrix. The average grain size of the 2
phase is 17 μm.
The results of strain-rate jump test are summarized in
Figure 2. It appears that stress follows a rather strong
dependency on temperature and strain rate. For iso-struc-
tural tests, (T = const.), the stress sensitivity parameter,
n, can be calculated as
n = ⎛⎝
⎜ ⎞
⎠
⎟∂
∂
ln
ln


(1)
The acquired data suggests that, according to Eq.1,
the overall results might be separated into two fields: a
low-stress regime with n = 3 (3.0 < n < 3.2) and an
upper, high-stress regime, with n = 5 up to 9. Contrary to
the low-stress field, where n is almost constant, in the
high-stress region n gradually increases with temperature
and stress. The boundary between low- and high-stress
regions may be approximated using a straight line and
denotes, in fact, the transition between different defor-
mation mechanisms.
The n = 3 value corresponds to alloy-type creep beha-
viour (Class A), the so-called viscous glide creep 15. In
contrast to this, metal-type creep behaviour (Class M) is
recognized with n = 5 and usually points to a dislocation
climb mechanism 16. However, the stress- dependent
gradual increase of n-values up to 9 suggests that power
law breakdown (PLB) may be more probable 17.
The apparent activation energy for creep is given by,
Q R
Tapp
= − ⎛⎝
⎜ ⎞
⎠
⎟∂
∂
ln
/


1
(2)
where R is the gas constant. For constant stress values in
the range of /μ = 4 × 10–4 – 2 × 10–3, the application of
Eq. 2 is illustrated on Figure 3. Calculated activation
S. TADI] et al.: DEFORMATION MECHANISMS IN Ti3Al-Nb ALLOY AT ELEVATED TEMPERATURES
358 Materiali in tehnologije / Materials and technology 44 (2010) 6, 357–361
Figure 3: Activation energy for deformation. (Filled symbols – low
stress regime in Figure 2 hollow – high stress regime.)
Slika 3: Aktivacijska energija derformacije (polni znaki – re`im
majhne napetosti na sliki 2, prazni znaki – re`im velike napetosti)
Figure 1: Microstructure of hot-rolled Ti3Al-Nb alloy; (OM, x200)
Slika 1: Mikrostruktura vro~e valjane zlitine Ti3Al-Nb (OM, pove~ava
200-kratna)
Figure 2: Strain-rates vs. normalized stress in the T = 1073–1373K
range
Slika 2: Odvisnost hitrost deformacje – normalizirana napetost v
obmo~ju T = 1073–1373 K
energies are shown in Figure 4. In the low-stress region,
activation energies are within the range typical of lattice
self-diffusion 10,11. In the high-stress regime, Q values
are higher and might be related to PLB 5. Stress-depen-
dent Q values require that, apart from the creep, dislo-
cation glide mechanisms are operative. However, it is
not yet clear which one is rate controlling.
Phenomenological relationship for power-law creep
takes the well known form




kT
D b
A
n
=
⎛
⎝
⎜
⎞
⎠
⎟ (3)
where A is a dimensionless constant, characteristic of
that particular creep mechanism, D is the effective
diffusion coefficient and all others have their usual
meaning.
The effective diffusion coefficient is calculated using
Eq. 4:
( )D D Q kT= −0 exp / (4)
where Q represents the mean calculated activation
energy and Do is the so-called frequency factor. For the 2
phase, Do = 2.24 × 10–5 m2/s and for ,, Do = 3.53 × 10–4
m2/s 18. The effective diffusion coefficient is calculated
to be within the range of 1.5 × 10–20 – 5 × 10–17 m2/s
which is in good agreement with 18. It should be pointed
out that these values correspond to volume self-
diffusion. For the viscous drag regime, appropriate
diffusion would be Darken’s chemical inter-diffusion
and, for lower temperatures, pipe (core) diffusion is
more likely to be operative 19. Anyway, due to the lack
of consistent published data, volume self-diffusion is
used in further calculations. By applying Eqs. 3 and 4 to
the experimental results, the plot of Figure 5 is revealed.
Two segments are clearly distinguishable: (i) a linear
one, with n = 3 and A 
 1 up to the normalized stress in
the vicinity of /μ = 10–3 and, beyond this stress, (ii) an
exponential upward curve with continuous slope n =
5–9. Stress /μ = 10–3 at the departure of linearity is
widely accepted as a characteristic for PLB. Figure 5
also includes theoretical calculations for viscous-glide
20 and climb-controlled creep 21. Reasonable accord is
more than clear, regardless of the diffusion approxima-
tions mentioned above.
It should be noted in Figure 5 that n = 5 is only a
small segment of the experimental data regarding higher
stresses. This fact supports some confirmations that the
behavior of n = 5 might be suppressed between viscous
glide and power-law-breakdown 22,23. In other words,
with increasing stress (or lowering temperature), defor-
mation advances sequentially: viscous-drag controlled
glide  climb + glide controlled flow.
It is obvious that Eq. 3, due to PLB, cannot describe
the overall deformation behaviour of the investigated
alloy. For this reason, a more flexible, hyperbolic sine
function is used:
S. TADI] et al.: DEFORMATION MECHANISMS IN Ti3Al-Nb ALLOY AT ELEVATED TEMPERATURES
Materiali in tehnologije / Materials and technology 44 (2010) 6, 357–361 359
Figure 5: Diffusion-compensated strain rate as a function of
normalized applied stress. Theoretical predictions of Eq. 3 are
included.
Slika 5: Hitrost deformacije kompenzirana z difuzijo v odvisnosti od
normalizirane delujo~e napetosti. Vklju~ene so teoreti~ne napovedi iz
ena~be. 3.
Figure 4: Semi-logarithmic stress- dependence of activation energies
Slika 4: Semilogaritmi~na odvisnost napetosti od aktivacijske energije
Figure 6: Hyperbolic sine fit (Eq. 5) of the data
Slika 6: Sinushiperboli~ni prikaz podatkov (ena~be 5)





kT
D b
A
n
=
⎛
⎝
⎜
⎞
⎠
⎟
⎡
⎣⎢
⎤
⎦⎥
sinh (5)
Eq. 5 gives a good description of the creep and
hot-working data, Figure 6. However, it is not the best
possible linear fit. It has been searched for reasonable
linear fit, yet parameters A,  and n reflect the physical
meaning of deformation. In Figure 6, stress exponent n
 3 reflects viscous drag at low stresses;  = 103 is the
inverse of normalized stress for PLB, i.e. sinh(/μ) = 1
and, A  10–9 points to diffusion-normalized strain-rate
for PLB. To repeat, some other values of A,  and n
might yield better linearity of Eq. 5. However, it would
be more valuable in technological modelling of hot-
working), than from the academic merit of identification
of deformation mechanisms.
4 DISCUSSION
Viscous glide has been reported to occur in a large
number of solid solution alloys. In particular, aluminium
alloys have been examined well 24–27. Also, this phe-
nomenon is observed commonly in a lot of intermetallics
28–31. Two mutually competitive mechanisms are opera-
tive in this stress regime, dislocation glide and climb.
Within the certain combination of materials’ and techno-
logical parameters glide is the slower one and thus
becomes the rate controlling mechanism of deformation.
Several possible viscous drag mechanisms were
proposed. Cottrell and Jaswon 32 proposed that the
dragging results from segregation of solute atmospheres
to moving dislocations. The dislocation speed is limited
by the rate of the solute atmosphere movement. Also,
segregation to stacking faults is proposed 33. Further,
stress induced local ordering might be an obstacle to
dislocation movement 34. Finally, in long-range ordered
alloys, as it is in the 2 phase, dislocations are impeded
due to anti-phase boundaries 35. If only solute atmo-
sphere dragging is considered, Mohamed and Langdon 24
proposed the relationship for creep in viscous glide:

 



=
− ⎛
⎝
⎜
⎞
⎠
⎟
( )1
6 2 5
3
kTD
e cb
(6)
where e is the solute-solvent size difference, c is the
concentration of solute atoms and D is the chemical
diffusion coefficient for the solute atoms. Eq. 6 serves
as an explanation that viscous glide occurs preferably in
materials with a relatively large atom size mismatch (e)
and higher solute concentrations (c). InTi3Al-Nb alloy,
both aluminium and niobium atoms act as solutes in the
-phase solid solution. In the 2 phase niobium may
substitute both Ti and Al atoms. The atomic volumes of
Ti and Nb are almost identical. However, the atomic
diameter of an Al atom is much lower than that of a
titanium atom 19. Also, viscous glide seems to be parti-
cularly strong at high temperatures where the volume
fraction of the -phase (Vf) dominates over 2. For
example, from the quasi-binary phase diagram Ti3Al-Nb
36, it can be calculated that Vf at T = 1373 K is 90 %, at
T = 1273 K Vf = 65 % while at T = 1073 K  is only 30
%. This suggests strongly that the aluminium-rich
-phase is predominantly in charge of viscous glide.
The transition from viscous glide to climb controlled
creep has been the subject of numerous investigations
37–41. There is a general agreement that the transition is
due to the breakaway of dislocations from solute atmo-
spheres. Under some critical stress the glide becomes
faster than the climb and the latter starts to be the rate
controlling mechanism. The break-away stress is propo-
sed 42:
 =
⎛
⎝
⎜
⎞
⎠
⎟
W
kTb
c
m
-12
2
3 (7)
Wm is the maximum interaction energy between a,
solute atoms and dislocations, c is the solute concen-
tration and  is the factor which depends on the nature of
the impurity clouds. Interaction energy Wm is
W Vam = −
+
−
⎛
⎝
⎜ ⎞
⎠
⎟1
2
1
1


 Δ (8)
Where  is Poisson’s ratio ( = 0.24),  is shear
modulus (Eq.1) and Va is volume difference between
solute and solvent atoms. From 19, VAl = 1.66 × 10–29 m3
and VTi = 1.81 × 10–29 m3, and Va = 0.15 × 10–29 m3.
With b = 2.83 × 10–10 m, and the atomic concentration of
aluminium cAl = 0.25, Eqs. 8 and 7 give good agreement
within the order of magnitude with the experimentally
determined /μ  10–3. However, there is no doubt that
interstitial impurities sustain break-away stress strongly.
While the viscous glide regime is clearly distin-
guishable, the distinction between five-power-law and
PLB isn’t so obvious. Close inspection of Figure 2, 5,
and 6 suggests strongly that the n = 5, regime is not
particularly a defined creep regime of this alloy, but
presents only the very beginning of PLB. Suppression of
five-power-law between viscous glide and PLB is also
found in some Al-Mg 43 and Al-Zn alloys 44.
Typically, within PLB, as n increases activation
energy decreases. However, in this work, a conflicting
result was found – activation energy is somewhat higher
as n increases. The behaviour of the Q values in Figure 4
excludes inaccurate scatter of data. Such a result might
be rationalized trough the fact that the ductile -phase
and more rigid 2-phase may behave quite differently
during deformation. Eq. 2, used to calculate activation
energy, is derived from
 = −⎛⎝
⎜ ⎞
⎠
⎟A
Q
RT
n exp (9)
If multiple deformation processes occurs through
sequential serial steps, (as glide and climb are), acti-
vation energy corresponds to the slowest one. However,
if parallel, simultaneous mechanisms are operative,
activation energy corresponds to the fastest one. Assum-
S. TADI] et al.: DEFORMATION MECHANISMS IN Ti3Al-Nb ALLOY AT ELEVATED TEMPERATURES
360 Materiali in tehnologije / Materials and technology 44 (2010) 6, 357–361
ing that applied stress is equally distributed on both
phases, the so-called iso-stress model 45 results in:
          = = ⇒ = +2 2 2V V (10)
meaning that the strain-rates of phases might depend
strongly on volume fractions (V). Combining Eqs 9, 10
yields
   

 

 = −
⎛
⎝
⎜
⎞
⎠
⎟ + −
⎛
⎝
⎜
⎞
⎠
⎟A V
Q
RT
A V
Q
RT
n n
2
2
2
2
exp exp (11)
A similar result can be obtained from a so-called
iso-strain-rate model 46 (   2 = ). The main con-
clusion from Eq.11 (and similar two-phase deformation
models 47–50), is that simultaneous mechanisms may be
operative in the 2 and  phases during deformation. In
that case, apparent activation energy might be reflected
by the 'faster' phase. The same issue might be seen in
Figure 5, where experimental results follow the faster of
the two theoretical mechanisms.
5 CONCLUSIONS
Strain-rate-jump tests were performed on two-phase
Ti3Al-Nb alloy in a wide range of temperatures
(1073–1373 K) and strain rates from (10–6 to 10–3 s–1).
Deformation mechanisms were identified in terms of
stress-exponent (n), activation energy (Q) and fitting
with known theoretical equations.
1. Viscous glide (n = 3 and Q close to lattice self diffu-
sion) was identified at a high temperature, low-stress
regime.
2. The five power law is suppressed between the vis-
cous glide and power law breakdown.
3. Power law breakdown begins at /μ  10–3, very
common in most metals and alloys, and features a
continuous, stress dependent, increase of n and Q.
4. Deformation behaviour of the alloy can be well
described with hyperbolic sine function.
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S. TADI] et al.: DEFORMATION MECHANISMS IN Ti3Al-Nb ALLOY AT ELEVATED TEMPERATURES
Materiali in tehnologije / Materials and technology 44 (2010) 6, 357–361 361

V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
SYNTHESIS AND CHARACTERISATION OF CLOSED
CELLS ALUMINIUM FOAMS CONTAINING DOLOMITE
POWDER AS FOAMING AGENT
PRIPRAVA IN KARAKTERIZACIJA ALUMINIJSKIH PEN Z
ZAPRTO POROZNOSTJO, IZDELANIH Z DOLOMITNIM PRAHOM
KOT SREDSTVOM ZA PENJENJE
Varu`an Kevorkijan1, Sre~o Davor [kapin2, Irena Paulin3, Borivoj [u{tar{i~3,
Monika Jenko4
1Independent Researcher, Betnavska cesta 6, 2000 Maribor, Slovenia
2Institut "Jo`ef Stefan", Jamova 39, 1000 Ljubljana, Slovenia
3Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
varuzan.kevokijan@impol.si
Prejem rokopisa – received: 2010-07-03; sprejem za objavo – accepted for publication: 2010-08-30
In this work, the viability of dolomite powder as cost-effective alternative to TiH2 foaming agent was investigated. Closed cells
aluminium foam samples were prepared starts from solid, foamable precursors synthesized by powder metallurgy and melt
route. Precursors obtained by melt route were machined and additional cold isostatic pressed in order to improve their density.
In all cases, the resulted precursors consisted of an aluminium matrix containing various mass fractions of uniformly dispersed
dolomite powders of various average particle size and 5 % of SiC particulates. Precursors were foamed by inserting into a
cylindrical stainless steel mould and placing inside a pre-heated batch furnace at 700 °C for 10 min.
The quality of foamable precursors was evaluated by determining their initial density and the foaming efficiency. On the other
side, the quality of the obtained foams were characterised by their density, microstructure and mechanical properties.
Experimental findings confirmed that aluminium foams synthesized with dolomite powder as blowing agent can be prepared by
both powder metallurgy and melt route, as well as that the density, microstructure, compression strength, and energy absorption
capacity are quite comparable with corresponding counterparts foamed by TiH2.
Key words: closed cells aluminium foams, dolomite particles as foaming agent, powder metallurgy, melt route, foaming
efficiency, mechanical properties
V delu preu~ujemo mo`nost nadome{~anja TiH2 kot sredstva za penjenje z dolomitnim prahom. Vzorce aluminijske pene smo
pripravili s prekurzorji, izdelanimi po postopku metalurgije prahov in po livarskem postopku. Prekurzorje, izdelane po
livarskem postopku, smo tudi strojno obdelali in nato {e hladno izostatsko stisnili, da smo pove~ali njihovo gostoto. Ne glede na
postopek njihove izdelave, so dobljeni prekurzorji vsebovali aluminijsko matriko in v njej razli~ne koncentracije enakomerno
porazdeljenih delcev dolomita ter SiC-delcev razli~nih povpre~nih velikosti. Pene smo v nadaljevanju izdelovali tako, da smo
prekurzor zaprli v za ta namen izdelan jekleni model, ki smo ga nato za 10 min vstavljali v pe~, ogreto na 700 °C.
Kakovost prekurzorjev za penjenje smo ugotavljali na osnovi njihove gostote in u~inkovitosti penjenje. Po drugi strani smo
kakovost izdelanih pen dolo~ali na osnovi njihove gostote, mikrostrukture in mehanskih lastnosti.
Eksperimentalne ugotovitve so potrdile, da je aluminijske pene mogo~e izdelovati z dolomitnim prahom kot sredstvom za
penjenje, tako po postopku pra{ne metalurgije kakor tudi po livarskem postopku in tako, da so njihova gostota, mikrostruktura,
tla~na trdnost in sposobnost absorpcije energije povsem primerljive z vrednostmi, ki jih navaja literatura za Al-pene, izdelane s
TiH2-penilom.
Klju~ne beside: aluminijske pene z zaprto poroznostjo, delci dolomita kot sredstva za penjenje, postopki metalurgije prahov,
livarski postopki, u~inkovitost penjenja, mehanske lastnosti
1 INTRODUCTION
Since last few decades, closed cells aluminium foams
– one of the lightest engineered materials have been
subject of investigation1,2 as candidate for a broad range
of applications. However, instead of significant progress
made in aluminium foams manufacturing, as well as
properties development and optimisation, a longed
expected commercialisation of this class of materials
remained poor. The limited commercial potential of
aluminium foams is mostly affected by its inadequate
viability (the insufficient balance between its technical
and economical attributes) caused by high cost and
properties which are not always in line with customers
expectations.
Closed cells aluminium foams are produced either by
(i) direct foaming methods starting from slurry of molten
aluminium or aluminium alloys and uniformly dispersed
non-metallic particles to which gas bubbles are added to
create foam, and (ii) indirect foaming methods starting
from a solid, foamable precursor which upon melting
expands and transforms to foam.3 The foamable pre-
cursor, aluminium-based matrix containing uniformly
dispersed blowing agent particles, can be produced either
by powder metallurgy4 or melt route.5
The direct foaming methods are cost-effectively but
result only in medium quality level of foams. On the
Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371 363
UDK 669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)363(2010)
other side, the closed cells aluminium foams produced
by indirect foaming of precursors made by powder
metallurgy route are with the superior quality but, at the
same time, very high cost. Precursors made by melt route
are with significantly lower cost, replacing the expensive
aluminium powder by conventional melt. The additional
reducing of cost can be achieved by replacing expensive
TiH2 with alternative inexpensive blowing agents, parti-
cularly carbonates such as CaCO3. The CaCO3 was
applied for successful indirect preparation of aluminium
foams by powder metallurgy and melt route.6 The usage
of other carbonates (e.g. MgCO3 and CaMg(CO3)2) was
also reported but not as often as CaCO3.
Among its significantly lower cost compared to the
cost of TiH2, CaCO3 as the blowing agent has also seve-
ral other advantages. As detailed discussed by Gergely et
al.,5 CaCO3 reacts with molten aluminium creating the
foaming gas (CO2) and various solid particles, depending
on the composition of the aluminium alloy (CaO, Al2O3,
Al4C3 and MgAl2O4). In contrast to TiH2 which decom-
position leads to the formation of chemically inert hydro-
gen, the CO2 foaming gas obtained by the decomposition
of CaCO3 reacts with melt stabilising the foam suspen-
sion. The results of Gergely et al.5 suggest that, as a
result of foaming gas (CO2)/melt reaction, a thin solid
reaction layer forms in the early stages of the foaming
process stabilising, by the beneficial effect on the surface
tension modification, the cells against coarsening and
coalescence. In addition, the solid particles obtained by
thermal decomposition of CaCO3 enhance the melt
viscosity, further promoting the stabilisation of the foam.
On the other side, the main disadvantage of CaCO3 as
foaming agent is its relatively high decomposition
temperature (between 700 °C and 900 °C) – significantly
above the melting point of pure aluminium and alumi-
nium alloys. High foaming temperature makes alumi-
nium foams stabilisation more demanding and costly.
However, an opposite problem exists with TiH2. It
decomposes far below the melting point of aluminium
and aluminium alloys and thereof should be surface
engineered8 in order to shift the decomposition tempera-
ture closed to the melting point of aluminium alloys,
which introduces additional cost.
One of the most important steps in production of
aluminium foams is their stabilisation. Either in direct or
indirect methods, foaming is always initiated by disper-
sing a large number of gas bubbles into a melt, leading to
the formation of stable slurry. The solid foam is then
fabricated by freezing the obtained slurry through solidi-
fication process. Thereof, during processing aluminium
foam is going through a series of transient states which
changes its morphology considerably. Generally, foams
are kinetically stable, if do not change significantly in
the time span between completion of the blowing pro-
cess and solidification. Although aluminium foam
stabilisation has been well discussed in literature and
numerous models of foam stability were presented,3 the
reason for their stability in the liquid or semi-liquid state
is still under dispute. However, it has been shown that
aluminium melts without a solid phase are not foamable.
Furthermore, it has been demonstrated that the stable
slurry of gas bubbles in molten metal is an important
prerequisite for obtaining more uniform microstructure
of the end product, without irregularities caused by
coarsening and coalescence. The effect of various
processing parameters (temperature, time, melt viscosity,
surface tension, wetting behaviour etc.) on the foama-
bility of aluminium alloys and stability of slurry of gas
bubbles in molten metal has been also investigated.4,9–11
It was found that aluminium foam slurry stabilisation is
the most effective at temperatures slightly below or
above the melting point of Al or the applied Al alloy.
The processing time should be properly controlled due to
the fact that the pore size and the total porosity volume
increase with time. Thereof, very short processing time
will result in fine pores but also the low total porosity
and low energy absorption capacity. The existence of the
solid particles in the melt is necessary for nucleation of
bubbles and the increase in viscosity. Viscosity is one of
the most important parameters influencing the foam
stability and should be kept inside the proper processing
window. An increase of viscosity generally enhances the
stability of the bubbles in slurry but also influences their
size and foam microstructure development. Finally,
lowering of the surface tension of the molten metal (e.g.
addition of magnesium) and improving wetting of solid
particles with melt significantly was found to enhance
significantly stabilisation of gas bubbles inside the
slurry.
Although the development of aluminium foams looks
back on a long history, none of the processes available
nowadays has been sophisticated to a level comparable
with that of polymeric foams. The reasons for that are
lack of understanding of the basic mechanism of alumi-
nium foaming, insufficient ability to make Al foams of a
constant quality, knowledge of aluminium foam proper-
ties is insufficient, physical properties of Al foams are
not good enough, transfer of research results to design
engineers is insufficient and Al foams are still too
expensive.3
Further improving of the competitiveness of the
indirect foaming of aluminium alloys it is particularly
important to formulate the cost-effective blowing agent
which thermal decomposition will proceed (would
happen) at temperature close to the melting point of the
selected alloy, evolving at the same time gaseous and
solid products well involved in stabilisation of foamed
slurry.
Considering that aluminium foam slurry stabilisation
is most effective at temperatures slightly below or above
the melting point of Al or the selected Al alloy, the most
suitable blowing agent should decomposes mostly inside
that temperature interval. Stabilisation of foam slurries at
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
364 Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371
temperatures far below or above the temperature of
liquidus is more difficult and costly.
In the present paper, the performance of dolomite
powder as a cost-effective foaming agent was investi-
gated in both powder metallurgy and melt route of foams
preparation. The influence of various processing para-
meters on foaming behaviour, as well as the development
of foam microstructure and mechanical properties was
also discussed.
2 EXPERIMENTAL PROCEDURE
All foams made in this work were prepared by
indirect foaming methods starts from a solid, foamable
precursor which consists of a metallic matrix containing
uniformly dispersed blowing agent particles. Foamable
precursors were made by: (i) powder metallurgy route
and (ii) melt route by using the same blowing agent –
dolomite powders (type A, B and C) with various
average particle sizes ((44, 76 and 97) μm, respectively).
The morphology of as received dolomite powders
was investigated by scanning electron microscopy
(SEM/EDS) whereas X-ray diffraction (XRD) measure-
ments were applied to identify the phases and their
crystal structure. In addition, the thermogravimetric
analysis (TGA) was also performed in a Setaram Labsys
DTA1600 equipment.
By following the powder metallurgy (P/M) route,
foamable precursors were made by mixing Al powder,
with an average particle size of 63 μm (purity: 99.7 %,
oxygen content: 0.25 %), 5 % of SiC particles with an
average particle size of 10 μm and (3, 5, 7 and 12) % of
blowing agent, followed by cold compaction, in a
lubricated 20 mm diameter die, to a pressure of 600MPa
to 900MPa.
In the case of melt route, foamable precursors with
the same concentration of dolomite blowing agent ((3, 5,
7 and 12) %) and the same geometry are prepared by a
induction heated batch-type stir-casted in which
aluminium powder (the same used for the P/M route)
was induction melted, followed by the addition of
dolomite particles, stirring and casting. Once molten
aluminium heated to 700 °C, the power addition was
switched off and melt stirring was initiated until the
temperature of melt decreased to 685 °C. After that, the
blowing agent/aluminium powder mixture (1 : 2 mass
ratio) was introduced and the melt was stirred (at
approximately 1200 r/min) for additional 30-90 s.
Finally, foamable precursors were prepared by casting
the semi-solid slurry into a room temperature mould with
20 mm diameter.
The solidified precursors were machined and some of
samples were additionally cold isostatically pressed.
The density of foamable precursors was calculated
from the mass and geometry of the samples and, in addi-
tion, measured by Archimedes method. The distribution
of blowing agent particulates inside Al matrix was
examined by analysing the optical and scanning electron
micrographs of as polished bars.
All precursors were foamed in a conventional batch
furnace with air atmosphere circulation under the same
experimental conditions (temperature, time, cooling
method). Before foaming, the individual precursors were
inserted into a cylindrical (40 mm in diameter, 70 mm
long) stainless steel mould coated with a boron nitride
suspension. The mould dimensions and the precursor
size (20 mm in diameter and 60 mm long) were selected
to allow an expansion of the precursor to foam with
theoretical density closed to 0.6 g/cm3. The arrangement
was placed inside a pre-heated batch furnace at 700 °C
for 10 min. After that period of time, the mould was
removed from the furnace and the foaming process was
stopped by rapid cooling with pressurised air to room
temperature. The thermal history of the foam sample was
recorded, using a thermocouple located directly in the
precursor material.
Foam density using the Archimedes method was
carried out. The porosity of manufactured foam was
calculated by the rate: 1-(foam density/aluminium den-
sity). Macro and microstructural examination was per-
formed on sectioned obtained by wire precision cutting
across the samples and on samples mounted in epoxy
resin, using optical and scanning electron microscopy
(SEM/EDS).
An average particle size of pores in the foams was
estimated by analysing the optical and scanning electron
micrographs of as polished foam bars using the point
counting method and image analysis and processing
software.
Regarding mechanical properties of the foams, the
uniaxial room temperature compressive tests were
carried out on a Instron 1255 testing machine at a con-
stant 5 mm/min crosshead displacement. Testing was
performed on standard prismatic foam specimens of 50
mm x 12 mm x 17 mm so that each point of the stress-
strain curve was determined as an average of four
individual measurements. Compression was stopped
whenever either 80 % strain or 95 kN force (equivalent
to 61.9 MPa) were reached. As a result of testing, the
uniaxial compression stress-strain curve, compressive
strength and energy absorption after a 30 % strain were
determined and correlated with the density, the average
pore size and microstructure of foam samples.
3 RESULTS AND DISCUSSION
3.1 Morphology investigation of performs and alumi-
nium foams
Morphology of the applied foaming agent – dolomite
powder grade A is presented in Figure 1. As evident,
dolomite particles are irregularly shaped and non-agglo-
merated.
XRD analysis of dolomite powder (grade A) revealed
the presence of about 5 % of calcium carbonate as an
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371 365
impurity, Figure 2. Thereof, the foaming agent (as
received dolomite powder) consisted of 95 % of dolo-
mite and 5 % of CaCO3.
Thermogravimetric (TGA) curve for the dolomite
powder is plotted in Figure 3. TGA analysis has shown
that dolomite powder undergoes thermal decomposition
above approximately 650 °C and that decomposition
ends about 830 °C. Because of that, higher foaming
temperatures are necessary than with TiH2 precursor,
particularly when higher foaming efficiency on final
foams is required. On the other side, higher onset
temperature of CO2 evolution from dolomite powder
enabling the incorporation of dolomite particles into
aluminium melt without the need of any special
pre-treatment in order to prevent premature gas-release.
Separate TGA measurement (not present here) showed
that the kinetics of gas-release (dolomite mass loss) are
slightly slower in the case of larger dolomite particulates.
Microstructures of cold isostatically pressed performs
obtained by powder metallurgy and melt route were also
analysed by SEM; Figures 4 and 5.
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
366 Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371
Figure 2: XRD of dolomite powder (grade A) showing the presence
of approx. 5 % of CaCO3 phase
Slika 2: XRD vzorca dolomitnega prahu (tip A), ki potrjuje prisotnost
pribl. 5 % CaCO3-faze
Figure 3: Thermogravimetric (TG) and differential scanning calori-
metric (DSC) measurement curves for dolomite powder grade A.
Slika 3: TGA dolomitnega prahu (tip A)
Figure 1: SEM micrograph showing the size and morphology of the
foaming agent – dolomite powder grade A.
Slika 1: SEM-posnetek dolomitnega prahu (tip A)
Figure 5: SEM micrograph of cold isostatically pressed perform
obtained by melt route
Slika 5: SEM-posnetek hladno izostatsko stisnjenega prekurzorja,
izdelanega po livarskem postopku
Figure 4: SEM micrograph of cold isostatically pressed preform
obtained by powder metallurgical route reveals almost fully dense
aluminium matrix, dolomite (dark) and silicon carbide (gray) particles
without particles agglomeration or particle fracture
Slika 4: SEM-posnetek hladno izostatsko stisnjenega prekurzorja,
izdelanega po postopku pra{ne metalurgije, ki razkriva gosto sintrano
aluminijsko matriko ter neaglomerirane delce dolomita (temno) in
silicijevega karbida (sivo)
The measured and calculated densities of foamable
performs obtained by P/M route, Table 1, confirmed that
under the applied pressure of isostatic pressing (740
MPa) all performs were with closed porosity and den-
sities above 98 % of theoretical. However, a density >99
% of theoretical was achieved for precursors containing
3–7 % of dolomite particles, with higher addition this
could not be achieved resulting in lower foaming
efficiency, as evident in Table 4. The foaming efficiency
of preforms was evaluated based on the relative density
of the obtained foam r, calculated by dividing the
apparent density of the foam F, with the density of the
aluminium, Al. Thereof, the foaming efficiency is
expressed as:
 = 1 – r = 1 – (F/Al) (1)
which actually corresponds to the volume fraction of
pores in foam samples. Lower the foam density, higher
is foaming efficiency.
Densities of as-machined foamable performs ob-
tained by melt route were significantly lower compared
to PM counterparts, Table 2. Moreover, in performs pre-
pared by melt route most of porosity was opened,
causing during foaming lower foaming efficiency, as
documented in Table 5. However, by additional isostatic
pressing, the remaining porosity in performs fabricated
by melt route was efficiently reduced below the volume
fraction 1.2 %, Table 3, enabling the formation of
aluminium foams with improved foaming performances,
Table 6.
In all cases, experimental results clearly indicate that
the porosity measured in foamable performs and the
apparent density achieved in aluminium foam samples
are inversely proportioned. Generally, foamable per-
forms with lower porosity resulted in foam samples with
higher apparent density and lower foaming efficiency.
Table 2: Porosity of as-machined foamable performs prepared by melt
route
Tabela 2: Poroznost strojno obdelanih prekurzorjev, izdelanih po
livarskem postopku.
Chemical composition of performs
w/%
Porosity
/%
Dolomite SiC Al powder Calculated Measured
Type-A
3 5 92 4.7 ± 0.47 4.9 ± 0.25
5 5 90 4.9 ± 0.49 5.1 ± 0.26
7 5 88 5.4 ± 0.54 5.8 ± 0.29
10 5 85 6.1 ± 0.61 6.6 ± 0.33
Type-B
3 5 92 4.5 ± 0.45 4.7 ± 0.24
5 5 90 4.7 ± 0.47 5.0 ± 0.25
7 5 88 5.0 ± 0.50 5.3 ± 0.27
10 5 85 5.7 ± 0.57 6.2 ± 0.31
Type-C
3 5 92 4.4 ± 0.44 4.5 ± 0.23
5 5 90 4.5 ± 0.45 4.7 ± 0.24
7 5 88 4.7 ± 0.47 5.0 ± 0.25
10 5 85 4.9 ± 0.49 5.3 ± 0.27
Table 3: Porosity of foamable performs obtained by melt route
improved by the additional isostatic pressing
Tabela 3: Poroznost prekurzorjev, izdelanih po livarskem postopku,
izbolj{ana z dodatnim hladnim izostatskim stiskanjem
Chemical composition of performs
w/%
Porosity
/%
Dolomite SiC Al powder Calculated Measured
Type-A
3 5 92 0.9 ± 0.09 0.9 ± 0.05
5 5 90 0.9 ± 0.09 0.9 ± 0.05
7 5 88 1.0 ± 0.10 1.1 ± 0.06
10 5 85 1.1 ± 0.11 1.2 ± 0.06
Type-B
3 5 92 0.9 ± 0.09 0.9 ± 0.05
5 5 90 0.9 ± 0.09 0.9 ± 0.05
7 5 88 0.9 ± 0.09 1.0 ± 0.05
10 5 85 1.0 ± 0.10 1.1 ± 0.06
Type-C
3 5 92 0.8 ± 0.08 0.8 ± 0.04
5 5 90 0.8 ± 0.08 0.8 ± 0.04
7 5 88 0.8 ± 0.08 0.9 ± 0.05
10 5 85 0.9 ± 0.09 1.0 ± 0.05
Table 4: Density, foaming efficiency and the average pore size of alu-
minium foams prepared by PM route
Tabela 4: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih po postopku pra{ne metalurgije
Initial composition of
foamable performs
w/%
Selected properties of foamed
samples
Dolomite SiC Alpowder
Density
(g/cm3)
Foaming
efficien-
cy (%)
Average
pore size
(mm)
Type-A
3 5 92 0.56 ± 0.03 79,3 0.9 ± 0.09
5 5 90 0.59 ± 0.03 78,1 0.9 ± 0.09
7 5 88 0.63 ± 0.03 76,7 1.1 ± 0.11
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371 367
Table 1: Porosity of foamable performs obtained by PM route
Tabela 1: Poroznost prekurzorjev izdelanih s postopkom pra{ne
metalurgije
Chemical composition of performs
w/%
Porosity
/%
Dolomite SiC Al powder Calculated Measured
Type-A
3 5 92 0.7 ± 0.07 0.7 ± 0.04
5 5 90 0.8 ± 0.08 0.8 ± 0.04
7 5 88 0.9 ± 0.09 1.0 ± 0.05
10 5 85 1.2 ± 0.12 1.3 ± 0.07
Type-B
3 5 92 1.0 ± 0.10 1.0 ± 0.05
5 5 90 1.0 ± 0.10 1.1 ± 0.06
7 5 88 1.2 ± 0.12 1.3 ± 0.07
10 5 85 1.6 ± 0.16 1.8 ± 0.09
Type-C
3 5 92 1.1 ± 0.11 1.1 ± 0.06
5 5 90 1.2 ± 0.12 1.3 ± 0.06
7 5 88 1.3 ± 0.13 1.4 ± 0.07
10 5 85 1.8 ± 0.18 2.0 ± 0.10
10 5 85 0.69 ± 0.03 74,4 1.3 ± 0.13
Type-B
3 5 92 0.51 ± 0.03 81,1 0.6 ± 0.06
5 5 90 0.53 ± 0.03 80,4 0.7 ± 0.07
7 5 88 0.57 ± 0.03 78,9 0.9 ± 0.09
10 5 85 0.59 ± 0.03 78,1 0.9 ± 0.09
Type-C
3 5 92 0.50 ± 0.03 81,5 0.6 ± 0.06
5 5 90 0.52 ± 0.03 80,7 0.6 ± 0.06
7 5 88 0.55 ± 0.03 79,6 0.8 ± 0.08
10 5 85 0.56 ± 0.03 79,3 0.9 ± 0.09
Table 5: Density, foaming efficiency and the average pore size of
aluminium foams prepared from as-machined foamable performs
fabricated by melt route
Tabela 5: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih iz strojno obdelanih prekurzorjev,
dobljenih po livarskem postopku
Initial composition of
foamable performs
w/%
Selected properties of foamed
samples
Dolomite SiC Alpowder
Density
(g/cm3)
Foaming
efficien-
cy (%)
Average
pore size
(mm)
Type-A
3 5 92 0.71 ± 0.04 73,7 1.2 ± 0.12
5 5 90 0.72 ± 0.04 73,3 1.4 ± 0.14
7 5 88 0.75 ± 0.04 72,2 1.4 ± 0.14
10 5 85 0.81 ± 0.04 70,0 1.5 ± 0.15
Type-B
3 5 92 0.61 ± 0.03 77,4 0.8 ± 0.08
5 5 90 0.63 ± 0.03 76,6 0.9 ± 0.09
7 5 88 0.66 ± 0.03 75,5 1.0 ± 0.10
10 5 85 0.70 ± 0.04 74,1 1.1 ± 0.11
Type-C
3 5 92 0.65 ± 0.03 75,9 0.8 ± 0.08
5 5 90 0.66 ± 0.03 75,5 0.9 ± 0.09
7 5 88 0.68 ± 0.03 74,8 1.1 ± 0.11
10 5 85 0.72 ± 0.04 73,3 1.5 ± 0.15
Table 6: Density, foaming efficiency and the average pore size of
aluminium foams prepared by melt route from as-machined and
additionally isostacially pressed foamable performs
Tabela 6: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih iz strojno obdelanih in hladno izostatsko
stisnjenih prekurzorjev, dobljenih po livarskem postopku
Initial composition of
foamable performs
w/%
Selected properties of foamed
samples
Dolomite SiC Alpowder
Density
(g/cm3)
Foaming
efficien-
cy (%)
Average
pore size
(mm)
Type-A
3 5 92 0.63 ± 0.03 76,7 1.1 ± 0.11
5 5 90 0.67 ± 0.03 75,2 1.2 ± 0.12
7 5 88 0.71 ± 0.03 73,7 1.3 ± 0.13
10 5 85 0.78 ± 0.03 71,1 1.6 ± 0.16
Type-B
3 5 92 0.57 ± 0.03 78,9 0.7 ± 0.07
5 5 90 0.59 ± 0.03 78,1 0.9 ± 0.09
7 5 88 0.62 ± 0.03 77,0 1.1 ± 0.11
10 5 85 0.63 ± 0.03 76,7 1.4 ± 0.14
Type-C
3 5 92 0.58 ± 0.03 78,5 0.7 ± 0.07
5 5 90 0.60 ± 0.03 77,8 0.8 ± 0.08
7 5 88 0.64 ± 0.03 76,3 1.0 ± 0.10
10 5 85 0.67 ± 0.03 75,2 1.1 ± 0.11
3.2 Microstructure investigation of aluminium foam
samples
The similar cellular structure development with
spherical, closed pores was obtained by both powder
metallurgy, Figure 6, and melt processing route, Figure
7. However, as evident in Figure 6, the samples obtained
by powder metallurgy route were with more uniform
microstructure consisted of well separated individual
cells. On the other hand, the microstructure of samples
obtained by melt processing route, Figure 7, revealed the
presence of non-uniformities caused by insufficient
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
368 Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371
Figure 7: Cross section of aluminium foam obtained by melting route
with characteristic channel network and foam drainage
Slika 7: Posnetek pre~nega prereza vzorca aluminijske pene, izdela-
nega po livarskem postopku, na katerem so poleg posameznih por
opazne tudi nehomogenosti, povzro~ene z zlitjem por in odtekanjem
taline skozi zna~ilne kanale v mikrostrukturi
Figure 6: Cross section of aluminium foams obtained by P/M route
with well separated individual cells and relatively uniform microstruc-
ture
Slika 6: Posnetek pre~nega prereza vzorca aluminijske pene, izdela-
nega po postopku pra{ne metalurgije, z izra`enimi posameznimi
porami in relativno enakomerno mikrostrukturo
stability of foams in transient states of their formation.
The most of observed imperfections were created by
flow (the movement of bubbles with respect to each
other), drainage (flow of liquid metal through the inter-
section of three foam films), coalescence (sudden insta-
bility in a foam film) and coarsening (slow diffusion of
gas from smaller bubbles to larger ones).
The absence of pore coarsening and drainage
suggests that there is a cell face stabilising mechanism
operating in the carbonate-foamed melts,5 slowing down
the cell face rupturing process and hence inhibiting cell
coarsening. The mechanism is likely to be a result of the
foaming gas (CO2)/melt or semi solid slurry reaction
during the foaming procedure which was detailed
discussed by Gergely et al.5
Concerning the average pore size and the uniformity
in cell size distribution, foams made by P/M route are
with finer pores and more regular morphology than
samples made by melt route, particularly these from
as-machined precursors. However, an additional cold
isostatically pressing of as-machined precursor obtained
by melt route was found to help in achieving more
uniform foams with smaller average pore size similar to
that obtained by P/M route. The improvement is most
probably caused by better compacting of individual
dolomite particles and aluminium matrix, resulting in
higher density of foamable precursor.
3.3 Mechanical properties
Figure 8 shows an example of stress-strain response
of samples foamed from performs prepared by P/M in
which the compressive strength of the foams was
correlated with their density.
Because of the closed cells structure, compressive
foam behaviour showed in all cases the typical stress-
strain diagram with a division on three parts: a linear
increase in strength mainly caused by elastic deforma-
tion, followed by a plateau caused by homogeneous
plastic deformation and a final step increase due to the
collapse of the cells. The compressive strength was taken
as the initial peak stress. Foams made by PM route
become with highest compressive strength, while
samples foamed from as-machined precursors were with
significantly lower values. For interval of foam densities
analysed in this work (from 0.5 g/cm3 to 0.7 g/cm3), it
was found that more dense foam samples shifted the
position of the plateau toward higher stress values.
The energy absorbed per unit volume (E-energy
absorption capacity), which is one of the most important
characteristic of aluminium foams, was determined by
the area under the stress-strain plots as follows:12
E
l
= ∫   ( )d0 (2)
where  is compressive stress, l is the limit of strain
concerned and  is compressive strain. The calculated
values of energy absorption capacity for samples were
plotted in Figure 9 and correlated with foams density.
The typical response was found to be a quasi-Gaussian
function with a maximum energy absorption capacity at
very narrow density range.
The maximum energy absorption capacity for various
foams is summarized in Figure 9. For foams made by
P/M route, the maximum energy absorption capacity of
6.82 MJ/m3 is achieved in foams with density of 0.63
g/cm3. On the other side, in samples foamed from
as-machined precursors fabricated by melt route, maxi-
mum energy absorption capacity of only 5.42 MJ/m3 was
detected. The maximum appeared at foam density of
0.65 g/cm3. Finally, in melt route fabricated precursors
additionally isostatically pressed before foaming, an
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371 369
Figure 9: Example of the optimization of aluminium foam density
range for maximum energy absorption capacity: A) foams obtained by
powder metallurgy, B) foams obtained from as-machined precursors
fabricated by melting route, and C) foams obtained from as-machined
and cold isostatically pressed precursors fabricated by melting route
Slika 9: Primer optimiziranja gostote vzorcev aluminijskih pen za
doseganje najve~je sposobnosti absorpcije energije: A) pene, izdelane
po postopku pra{ne metalurgije; B) pene iz strojno obdelanih preku-
zorjev, narejenih po livarskem postopku; C) pene iz strojno obdelanih
in hladno izostatsko stisnjenih prekurzorjev, narejenih po livarskem
postopku
Figure 8: The stress-strain response of various aluminium foam
samples from performs obtained by PM
Slika 8: Krivulja napetost – deformacija za vzorce aluminijskih pen
na osnovi predoblik, izdelanih po postopku pra{ne metalurgije
intermediate maximum energy absorption capacity of
6.23 MJ/m3 was detected in samples with density of 0.63
g/cm3.
The foaming process principally does not affect the
properties of the cell-wall material. However, it leads to
a unique spatial distribution of aluminium which results
in significantly different properties of foamed compo-
nent in comparison with a bulk part. It is obvious that the
properties of aluminium foam significantly depend on its
porosity, so that a desired property (or combination of
properties) can be tailored by selecting the foam density.
The mechanical properties of foams obtained by
applying dolomite powder as foaming agent are fully
comparable with counterpart properties in foams fabri-
cated by using TiH2.
4 CONCLUSION
The following conclusions can be drawn from this
work.
TiH2 powder as foaming agent was successfully
replaced by commercial dolomite powders of different
average particle size, fabricated by milling of nature
mineral.
Foaming precursors with different volume portions
(3–10 %) of dolomite powder particles as blowing agent
were routinely prepared neither by powder metallurgy or
melt route. Precursors obtained by powder metallurgy
were with superior homogeneity and densities above 98
% of theoretical. The counterparts obtained by melt route
were with more agglomerated dolomite particles and
significantly lower densities (93.4 % of theoretical).
However, it was found that an additional cold isostatic
pressing of these precursors improved their densities up
to 97.4 % of theoretical.
Density above 99 % of theoretical was achieved only
for precursors obtained by PM and melt route (improved
by additional isostatic pressing), containing 3-7 % of
dolomite particles of an average particle size of 44 μm.
With higher addition of dolomite particles and by using
dolomite powders with higher average particle size (76
μm and 97 μm) densities 99 % of theoretical could not
be achieved.
The foaming efficiency of experimentally prepared
precursors was evaluated based on the relative density of
the obtained foams (the apparent density of the foam
divided by the density of aluminium). The experimental
findings shown that the apparent density of foam
samples is inversely proportional to the density of
foaming precursor. Thereof, foamable precursors with
higher density resulted in foam samples with lower
apparent density and higher foaming efficiency. Foaming
of almost fully dense precursors (densified above 99 %
of theoretical density) resulted in foam samples with the
lowest densities (typically from 0.50 g/cm3 to 0.6 g/cm3)
and hence high foaming efficiency (75 %). In opposite,
foaming of precursors densified below 97 % of theore-
tical led to foams with higher densities (from 0.7 g/cm3
to 0.8 g/cm3) and lower foaming efficiency (below 75
%).
Under the same foaming conditions (temperature,
time), the average pore size of foam samples is
influenced by the density of the foaming precursors and
the initial size of the foaming particles. As a rule, in
foams made from precursors with high density (99 %
of theoretical), the average pore size remained below 1
mm. On the other side, in foams made from precursors
with lower density (below 97 % of theoretical), pores
grown to 20 % to 50 % higher average pore size.
Regarding the initial size of the foaming particles, which
also influences the density of precursor and hence the
density of the foam samples, an increase of the average
particle size of dolomite foaming agent was observed to
have the detrimental influence on the average size of
pores. Coarser dolomite powders led to the formation of
larger bubbles in foam structure. By introducing the term
of foaming efficiency, which in simple and transparent
way includes all experimental influences, it was possible
to establish the direct correlation between the foaming
efficiency and the average pore size as one of the main
parameters of the microstructure of foamed samples. The
experimental findings confirmed that the foaming
efficiency and the average pore size of foaming samples
are generally reciprocally depended. Thereof, higher
foaming efficiency results in foam microstructure with
finer pores.
Mechanical properties (compression strength and
energy absorption capacity) of foamed samples are also
strongly influenced by foaming efficiency. For interval of
foam densities analysed, the compression strength,
considered as the initial peak stress, was found to be
superior (approx. 13.8 MPa) in samples with increased
density (0.69 g/cm3) and thereof lower foaming effi-
ciency (74.4 %). In opposite to that, the maximum
energy absorption capacity was achieved in foams with
the highest foaming efficiency. The stress-strain response
of foamed samples consisted of three parts: a linear
increase in strength mainly caused by elastic defor-
mation, followed by a plateau caused by homogeneous
plastic deformation and a final step increase due to the
collapse of the cells. On the contrary, the energy
absorption capacity is a quasi-Gaussian function of
density, approaching maximum absorption capacity at
very narrow density range.
From experimental findings is obvious that the
properties of aluminium foam significantly depend on its
porosity, so that a desired property (or combination of
properties) can be tailored by selecting the foam density.
Foams synthesized with dolomite powder particles as
blowing agent can be cost effectively prepared by both
powder metallurgy and melt route. However, the decom-
position characteristics of dolomite powder enabling the
foaming of aluminium and aluminium alloys only above
its melting point.
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
370 Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371
Experimental findings confirm that microstructure,
compression strength and energy absorption capacity of
aluminium foams prepared by dolomite powder as
foaming agent are quite comparable with these in coun-
terparts foamed by TiH2.
5 REFERENCES
1 J. Banhart, Int. J. Vehicle Design, 37 (2005) 2/3, 114
2 E. M. A. Maine, M. F. Ashby, Mater Design, 23 (2002) 3, 307
3 J. Banhart, Adv. Eng. Mater., 8 (2006) 9, 781
4 S. Asavavisithchai, A. R. Kennedy, Scripta Mater., 54 (2006), 1331
5 V. Gergely, D. C. Curra, T. W. Clyne, Composites Science and
Technology, 63 (2003), 2301
6 L. E. G. Cambronero, J. M. Ruiz-Roman, F. A. Corpas, J. M. Ruiz
Prieto, Journal of Materials Processing Technology, 209 (2009),
1803
7 US Pat. No. 7.452.402, issued November 18, 2008
8 B. Matijasevic, J. Banhart, Scripta Mater., 54 (2006), 503
9 G. Kaptay, Colloids and Surfaces A: Physicochemi. Eng. Aspects,
230 (2004), 67
10 V. C. Srivastava, K. L. Sahoo, Materials Science – Poland, 25 (2007)
3, 733
11 C. C. Yang, H. Nakae, Journal of Materials Processing Technology,
141 (2003) 202
12 S. Asavavisithchai, R. Tantisiriphaiboon, Chiang Mai J. Sci., 36
(2009) 3, 302
Acknowledgement
This work was supported by funding from the Public
Agency for Research and Development of the Republic
of Slovenia, as well as the Impol Aluminium Company
and Bistral, d. o. o., from Slovenska Bistrica under con-
tract No. 2410-0206-09.
V. KEVORKIJAN ET AL.: SYNTHESIS AND CHARACTERISATION OF CLOSED CELLS ALUMINIUM FOAMS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 363–371 371

N. ORE[EK ET AL.: IMOBILIZACIJA LETE^EGA PEPELA Z ZASTEKLITVIJO
IMOBILIZACIJA LETE^EGA PEPELA Z ZASTEKLITVIJO
FLY ASH IMMOBILIZATION WITH VITRIFICATION
Natalija Ore{ek1, Franc Berk2, Niko Samec1, Franc Zupani~1
1Univerza v Mariboru, Fakulteta za strojni{tvo, Smetanova 17, 2000 Maribor, Slovenija
2Javno podjetje za komunalne storitve Roga{ka Slatina, d. o. o., Celjska cesta 12, 3250 Roga{ka Slatina, Slovenija
natalija.oresek@uni-mb.si
Prejem rokopisa – received: 2010-05-17; sprejem za objavo – accepted for publication: 2010-09-03
Raziskali smo mo`nost imobilizacije lete~ega pepela iz se`iga komunalnih odpadkov. Lete~i pepel smo skupaj s slovenskim
odpadnim steklom stalili in ulili v kovinske kokile, da je nastal prete`no amorfen produkt, tj. steklo, ki smo ga toplotno obdelali,
da se je tvorila steklokeramika.
Lete~i pepel je bil sestavljen iz delcev velikosti <1 μm, ki se povezujejo v aglomerate velikosti okrog 180 μm. Tali se pri
temperaturi okoli 1300 °C in se v teko~em stanju dobro me{a z odpadnim steklom. Steklo, nastalo pri litju, se pretvori v
steklokeramiko pri `arjenju v temperaturnem obmo~ju med 700 °C in 950 °C, pri ~emer se tla~na trdnost pove~a za dvakrat.
Vrednosti izlu`evanja te`kih kovin iz vzor~nega lete~ega pepela prekora~ujejo predpisane mejne vrednosti za odlaganje na
odlagali{~ih za nevarne odpadke, medtem ko so vrednosti za nastalo steklo in steklokeramiko pod predpisanimi vrednostmi.
Zasteklen lete~i pepel lahko odlo`imo kot stabiliziran in nereaktiven odpadek na odlagali{~ih nenevarnih odpadkov, mogo~e pa
bi ga bilo tudi predelati in uporabiti v koristne namene.
Klju~ne besede: lete~i pepel, zasteklitev, devitrifikacija, mehanske lastnosti, izlu`evanje
The possibility of fly ash immobilization was investigated. Municipal solid waste fly ash mixed with waste glass cullet was
melted and poured into metal moulds. Mainly amorphous glass product was formed, which was then heat treated to obtain glass-
ceramics.
Fly ash consisted of particles with size <1 μm, normally joined to larger agglomerates (180 μm). Its melting temperature was
around 1300 °C and it mixed well with the glass cullet in the liquid state. Glass formed during solidification was transformed
into the glass-ceramics after heat treatment in the temperature range between 700 °C in 950 °C, that doubled the compressive
strength. Leaching of heavy metals from fly ash exceeded the allowed values for dumping dangerous waste in landfill, whereas
the values for the produced glass and glass-ceramics were under the allowed values. The produced glass and glass-ceramics can
be landfilled as a stabilised and unreactive waste in landfills for nonhazardous wastes. Their properties also offer the possibility
for manufacturing useful products.
Key words: fly ash, vitrification, devitrification, mechanical properties, leaching
1 UVOD
Dejstvo je, da vsaka dejavnost povzro~a odpadke,
prav tako pa imajo tudi sodobne tehnologije predelave in
recikliranja odpadkov svoje stranske produkte v obliki
odpadkov 1. ^eprav si moramo prizadevati za nenehno
zmanj{evanje odpadkov na izviru njihovega nastanka ter
za pove~evanje stopnje recikla`e in snovne izrabe
posameznih frakcij, je del ostanka odpadkov smiselno
(in po direktivi EU tudi nujno) termi~no obdelati in izko-
ristiti njihovo energijsko vrednost. Neizogibni produkt
termi~ne obdelave odpadkov (se`iga komunalnih odpad-
kov) je trdni preostanek po se`igu. Odvisno od tehnike
se`iga od 10 % do 25 % 2 mase vhodnega materiala iz
se`iga izide kot trdni preostanek (pepel z re{etke, pepel
iz kotla in iz odpra{evalnih naprav (elektrofiltrski lete~i
pepel) in filtrna poga~a). Njihova sestava je zelo razli~na
in odvisna predvsem od sestave odpadkov, ki jih se`i-
gamo, in tehnike zgorevanja. Vnaprej je ne moremo
napovedati, saj je odvisna od ve~ dejavnikov, kot so letni
~as, spremembe obmo~ja pobiranja odpadkov … Zaradi
vedno ve~jih potreb po se`iganju, bodo vedno ve~je tudi
potrebe po ve~jih (varnih) odlagali{~ih, zato je potrebno
razviti postopke, ki bodo tudi nevarne trdne preostanke
po se`igu imobilizirali, da bodo ~im manj nevarni ljudem
in okolju.
Lete~i pepel je generi~ni izraz za vse tipe drobno-
pra{natega pepela, ki se zbira v sistemu ~i{~enja dimnih
plinov in se iz toka dimnih plinov odstrani, {e preden se
dimnim plinom dodajo sorbenti. Med potovanjem
dimnih plinov do ~istilne naprave se na drobne delce
ogljika v lete~em pepelu adsorbirajo nevarne kovine (Cd,
Cr, Cu, Hg, Pb, …), kot tudi dioksini ter furani in drugi
kondenzibilni ogljikovodiki, zato ima lete~i pepel zaradi
prese`enih mejnih vrednosti izlu`kov posameznih
toksi~nih komponent zna~ilnosti nevarnega odpadka. V
skladu z zakonodajo 3 moramo odlagati lete~i pepel na
odlagali{~a za nevarne odpadke. Teh v Sloveniji ni, zato
je mogo~ samo izvoz teh odpadkov v tujino, kjer jih
deponirajo v zaprtih podzemnih prostorih iztro{enih
rudnikov premoga in rudnin (npr. KCl) 4.
Sestava lete~ega pepela je ve~inoma podobna tisti, ki
se uporablja v steklarski industriji pri proizvodnji stekla,
saj vsebuje 5: 11–35 % SiO2, 5–19 % Al2O3, 19–29 %
CaO (to so masni dele`i) in razli~ne koli~ine MgO,
Na2O, K2O in `vepla. Tako lahko glede na primerno
koli~ino steklotvorcev v lete~em pepelu dobimo po
taljenju in ohlajanju prakti~no inerten stekleni produkt, v
Materiali in tehnologije / Materials and technology 44 (2010) 6, 373–378 373
UDK 628.477:666.1.038 ISSN 1580-2949
Izvirni znanstveni ~lanek/Original scientific article MTAEC9, 44(6)373(2010)
katerega so vgrajene te`ke kovine in druge nevarne
snovi.
Taljenje lete~ega pepela po navadi poteka pri tem-
peraturah vi{jih od 1300 °C (5,6,7,8). Med taljenjem se
sestava pepela spremeni v odvisnosti od vsebnosti
anorganskih komponent, kajti komponente z ni`jim
vreli{~em delno izparijo (Cd, Pb, NaCl, KCl …) 9, zato
se posledi~no zmanj{a tudi masa preostanka. Pri tem
nastane iz pra{natega lete~ega pepela z majhno nasipno
gostoto kompaktno steklo, ki ima mnogo ve~jo gostoto
in s tem manj{o prostornino. Dosedanje raziskave 4–18 so
pokazale, da je treba za doseganje ustreznih mehanskih
in kemijskih lastnosti vzpostaviti pravo razmerje med
Si/Al/Ca v proizvodu. Zato je v nekaterih primerih treba
lete~emu pepelu dodati steklotvorne komponente, kot so
SiO2, TiO2 in MgO, da dobimo proizvod, ki ustreza
razredu diopsidov10, ali pa steklene ~repinje 6,11, ki vse-
bujejo veliko SiO2 in CaO, ki sta steklotvorna in
pospe{ita rast kristalov v kon~nem produktu. Ker je v
Sloveniji veliko odpadnega stekla (komunalni odpadki iz
gospodinjstev vsebujejo okrog 7 % stekla, komunalni
odpadki iz industrije 9 %, na razpolago pa je {e veliko
odpadnega embala`nega stekla, kar je skupaj okoli
50 000–60 000 ton odpadnega stekla), je zelo smotrno,
da se lete~emu pepelu doda odpadno steklo, saj potem
tudi to pridobi dodano vrednost.
S kasnej{o naknadno toplotno obdelavo v vzorcu
zasteklenega produkta spro`imo usmerjeno kristalizacijo,
razsteklitev oz. devitrifikacijo in s tem nastanek steklo-
keramike, ki ima bolj{e mehanske lastnosti.
Glavni cilj raziskave je bil ugotoviti, ali lahko z do-
datkom slovenskega odpadnega stekla k lete~emu pepelu
pridobimo trden in kemijsko obstojen produkt, ki bi ga
bilo mogo~e uporabiti za koristne izdelke (nasipni ma-
terial, polnilo za asfalt, tlakovci, bloki za gradnjo, zvo~ne
in zadr`evalne pregrade na cestah, steklena vlakna –
izolacije, stre{niki ...) ali vsaj za stabilno, okolju ~im
manj {kodljivo in dolgoro~no nadzorovano odlaganje,
skladno s slovensko zakonodajo 3.
2 EKSPERIMENTALNO DELO
Lete~i pepel po se`igu komunalnih odpadkov, skupaj
s kotlovskim prahom, ki se nabira v kotlovskih ceveh,
smo pridobili iz objekta za termi~no obdelavo komunal-
nih odpadkov AVA Abfallverwertung Augsburg GmbH,
Nem~ija. Odpadno komunalno steklo smo dobili od
podjetja DINOS (masni dele`i sestavin v odstotkih: 71 %
SiO2, 2 % Al2O3, 10 % CaO, 3 % MgO, 13 % Na2O +
K2O), industrijsko steklo pa iz podjetja Steklarna Ro-
ga{ka, d. d. (masni dele`i sestavin: 72,8 % SiO2, 4,41 %
CaO, 8.02 % N2O, 10,8 % K2O, 0,28 % Pb, 2,36 % BaO,
0,52 % Sb2O3, 0,51 % B2O3, …).
Granulometrijsko sestavo lete~ega pepela in s tem
povpre~no velikost delcev smo izmerili z aparaturo za
lasersko difrakcijo velikosti delcev Fritsch Analysette
22, njegovo nasipno gostoto pa v skladu s standardom
ISO 903. Mikrostrukturo vzorcev smo raziskali v svet-
lobnem mikroskopu Nikon Epiphot 300 in v vrsti~nem
elektronskem mikroskopu Sirion 400 NC, FEI Company,
opremljenim z EDS-analizatorjem INCA 350, Oxford
Instruments. Rentgenska fazna analiza je bila izvedena z
dvodimenzionalno 2D XRD-analizo (sinhrotronski rent-
genski `arki valovne dol`ine 0,1 nm; Elettra, Sincrotrone
Trst, Italija).
Pri preiskavi smo uporabili dva vzorca:
• 25 % lete~ega pepela in 75 % odpadnega stekla (od
tega je bilo 2/3 komunalnega embala`nega stekla in
1/3 industrijskega navadnega stekla). Oznaka vzorca
P25.
• 50 % lete~ega pepela in 50 % odpadnega stekla (ena-
ko razmerje kot pri P25). Oznaka vzorca P50.
Diferen~no vrsti~no analizo DSC 19 in termogravi-
metri~no analizo (TG) lete~ega pepela smo izvedli v
napravi STA 449 Jupiter podjetja Netzsch s hitrostjo
segrevanja in ohlajanja 20 °C/min. Z njo smo ugotovili
temperaturo taljenja in gravimetri~ne spremembe v masi
lete~ega pepela.
Lete~i pepel (pribli`no 25 g) smo talili pri tem-
peraturah 1290 °C in 1390 °C v korundnem lon~ku. Po
20 min zadr`evanja smo talino ulili v segret steklarski
model, po strditvi smo vzorce popu{~ali dobro uro pri
450 °C, da smo prepre~ili nastanek notranjih napetosti in
s tem razpad vzorcev. Tvorila se je steklu podobna snov
temno zelene, skoraj ~rne barve.
Vzorce smo nato preiskali z DSC-analizo, da bi ugo-
tovili optimalne razmere za pretvorbo stekla v steklo-
keramiko.
Zasteklene vzorce smo nato toplotno obdelali
(devitrificirali) v dveh stopnjah z `arjenjem v cevni pe~i:
najprej 1 h pri temperaturi 700 °C in nato 2 h pri 950 °C.
Za merjenje trdote vzorcev po Vickersu smo izbrali
obmo~je majhnih obremenitev med 200 g (oznaka HV
0,2) in 400 g (oznaka HV 0,4). Tla~no trdnost smo
izmerili na napravi Instron 1255.
Izlu`evanje je potekalo v skladu s standardom SIST
EN 12457-4 (24-urno izlu`evanje z vodo; razmerje
voda/trdna snov = 10 : 1), merjenje lu`ilnih parametrov
pa z metodo ICP-MS in ionsko kromatografijo.
3 REZULTATI IN DISKUSIJA
3.1 Analiza lete~ega pepela
Z lasersko difrakcijo smo ugotovili, da so delci
lete~ega pepela veliki od 1 μm do ve~ kot 200 μm,
vendar je bilo najve~ delcev v velikostnem razredu
100–200 μm. Iz tega sklepamo, da se delci povezujejo v
aglomerate, kar potrjuje tudi elektronski mikroposnetek
(slika 1).
Kvalitativna EDS-analiza vzorca lete~ega pepela je
pokazala, da je sestavljen prete`no iz O, Ca, Cl, Na, Si in
Fe. Difraktogram lete~ega pepela (slika 2, spodnja
krivulja) je imel 40–60 vrhov. Med njimi prevladujejo
vrhovi, ki pripadajo mineralom halita (NaCl), silvita
N. ORE[EK ET AL.: IMOBILIZACIJA LETE^EGA PEPELA Z ZASTEKLITVIJO
374 Materiali in tehnologije / Materials and technology 44 (2010) 6, 373–378
(KCl), -kremena (-SiO2), kalcijevega sulfata (CaSO4),
kalcijevega oksida (CaO) in hematita Fe2O3.
Simultana termi~na analiza je pokazala, da se je
taljenje lete~ega pepela pri segrevanju s hitrostjo 20
°C/min za~elo pri temperaturi 1197 °C, povsem pa se je
stalilo pri temperaturi 1230 °C. Pri segrevanju in
ohlajanju je bila izguba mase 35,6 %. Na osnovi teh
rezultatov sta bili izbrani temperaturi taljenja 1290 °C in
1390 °C. Ustrezni rezultati so bili dose`eni `e pri ni`ji
temperaturi 1290 °C, kar pomeni, da lahko prihranimo
pri energiji, pa tudi masa vzorca se je zaradi kraj{ega
zadr`evanja pri visokih temperaturah zmanj{ala le za
14,9 %, kar je bistveno manj kot pri simultani termi~ni
analizi.
3.2 Analiza stekel po taljenju
@e brez dodatka odpadnega stekla kot steklotvornega
faktorja, pridobimo z visokotemperaturnim taljenjem iz
pra{natega lete~ega pepela kompaktno steklo (slika 3a)
20, v na{i raziskavi pa smo `eleli pokazati, da se z do-
datkom stekla pomembno spremenijo fizikalne lastnosti
(trdota in tla~na trdnost) nastalega produkta. Pregled s
svetlobnim mikroskopom je pokazal, da po taljenju in
strjevanju nastanejo pore, vendar je produkt v splo{nem
kompakten in nekru{ljiv. Preiskava z vrsti~nim elek-
tronskim mikroskopom (sliki 3b in 3c) je pokazala, da je
mikrostruktura enakomerna, v stekleni osnovi so svet-
lej{i delci, katerih dele` se z ve~anjem dodatka odpad-
nega stekla zmanj{uje. Z EDS-analizo smo ugotovili, da
je v vzorcu s 25 % lete~ega pepela (P25) s sestavo
N. ORE[EK ET AL.: IMOBILIZACIJA LETE^EGA PEPELA Z ZASTEKLITVIJO
Materiali in tehnologije / Materials and technology 44 (2010) 6, 373–378 375
Tabela 1: Srednje vrednosti trdote in tla~ne trdnosti vzorcev nastalega stekla in steklokeramike
Table 1: Average Vickers hardnesses and compressive strengths of investigated samples
vzorec
sample
trdota po Vickersu HV
Vickers hardness
tla~na trdnost /MPa
compressive strength /MPa
steklo P25 (25 % lete~ega pepela)
glass (25 % fly ash) 665 HV 0,4 55
steklo P50 (50 % lete~ega pepela)
glass (50 % fly ash) 718 HV 0,4 42
steklokeramika (25 % lete~ega pepela)
glass-ceramics (25 % fly ash) 715 HV 0,4 105
steklokeramika (50 % lete~ega pepela)
glass-ceramics (50 % fly ash) 608 HV 0,4 148
Tabela 2: Rezultati izlu`evanja za izbrane elemente; vsi rezultati so v mg/kg s.s.
Table 2: The results of leaching for selected elements; all results are in mg/kg s.s.
element
lete~i
pepel
nastalo
steklo
nastalo
steklo
Steklo
keramika
iz P25
Steklo
keramika
iz P50
Zakonodaja 3: Mejne
vrednosti parametrov izlu`ka
stabiliziranih in nereaktivnih
nenevarnih odpadkov
Zakonodaja 3: Mejne
vrednosti parametrov
izlu`ka nevarnih
odpadkov
fly ash
formed
glass
P25
formed
glass
P50
Glass-
ceramics
P25
Glass-
ceramics
P50
(Slovenian legislation 3:
Leaching parameters for
stabilized and unreactive non
hazardous waste)
(Slovenian legislation 3:
Leaching parameters for
hazardous waste)
Pb 22,40 0,01 0,03 <0,01 <0,01 10 50
Ba 4,76 0,1 2 1,1 0,5 100 300
Cd 0,007 0,018 0,02 <0,001 0,001 1 5
Cu 4,27 0,05 0,07 0,04 0,03 50 100
Se 1,140 <0,01 <0,01 <0,01 <0,01 0,5 7
Sb 1 0,04 0,1 0,04 0,1 0,7 5
Cr 26,7 < 0,01 < 0,01 < 0,01 < 0,01 10 70
As 0,292 <0,01 <0,01 <0,01 <0,01 2 25
Kloridi
chlorides (Cl–) 128100 10 10 10 10 15000 25000
Slika 1: Delec lete~ega pepela z velikostjo okoli 180 μm je aglomerat,
sestavljen iz manj{ih delcev (SEM, sekundarni elektroni)
Figure 1: The particle of fly-ash, with the size of approximately 180
μm is an agglomerate consisting of several micron-sized particles
(SEM, secondary electron image)
osnove pribli`no Al0,01Ca0,04Cr0,04K0,02Mg0,02Na0,06Si0,25
O0,52 ter {e nekaj desetink Cl, Ti, Fe, Co, Ni, Zn in Sb,
medtem ko je v vzorcu s 50 % lete~ega pepela (P50)
sestava osnove: Al0,02Ca0,06Cr0,04K0,014Mg0,01Na0,04Si0,20
O0,55 in podobne koli~ine drugih elementov kot pri
prej{njem vzorcu. V vzorcu z manj pepela je ve~ K in
Na, kar ka`e, da so v tem primeru manj{e izgube
lete~ega pepela. V mikrostrukturi so tudi ostrorobi
svetlej{i delci, v vzorcu ~istega lete~ega pepela je to
predvsem spinel (Mg,Zn,Fe)O × (Al,Cr,Fe)2O3. V naj-
svetlej{ih delcih, ki so bolj ali manj okrogle oblike, so
te`ki elementi, kot sta npr. Sb in Pb. Dele` kristalnih faz
v steklu je zelo majhen, saj je na difraktogramu (srednja
krivulja na sliki 2) opazen zelo {irok vrh, ki je zna~ilen
za amorfno stanje, medtem ko ni opaziti uklonskih vrhov
kristalnih faz.
Izmerjena trdota nastalih stekel je med 650 HV 0,4 in
720 HV 0,4 (tabela 1), kar je precej ve~ od trdote Na-Ca
stekel (550 HV). Trdoti obeh stekel se med seboj
razlikujeta za pribli`no 50 enot, vi{ja trdota je bila
izmerjena pri vzorcu z ve~jim dele`em odpadnega stekla.
Pri preizkusih je bilo opaziti ve~ji raztros rezultatov, kar
je pri~akovano za tako krhek material. Podobno je bilo
tudi pri tla~nih preizkusih, pri ~emer sta bili trdnosti
obeh vrst stekel primerljivi (tabela 1).
Rezultati izlu`evanja (tabela 2) poka`ejo, da smo z
zasteklitvijo lete~ega pepela z odpadnim embala`nim
Na-Ca steklom pridobili steklene produkte, katerih
izlu`evanje kemikalij z vodo je dale~ pod nivojem vred-
nosti predpisanih parametrov, ki dovoljujejo odlaganje
na odlagali{~ih za stabilizirane in nereaktivne nenevarne
odpadke 3.
N. ORE[EK ET AL.: IMOBILIZACIJA LETE^EGA PEPELA Z ZASTEKLITVIJO
376 Materiali in tehnologije / Materials and technology 44 (2010) 6, 373–378
Slika 4: DSC-krivulja segrevanja nastalega stekla za vzorec s 50 %
lete~ega pepela (P50) (hitrost segrevanja je bila 20 °C/min)
Figure 4: DSC-trace of a produced glass, a sample with 50 % fly ash
(P50) (the heating rate was 20 °C/min)
Slika 3: Elektronski mikroposnetki vzorcev stekla po zasteklitvi: a)
lete~i pepel brez dodatka odpadnega stekla, b) vzorec P50 s 50 % lete-
~ega pepela c) vzorec P25 s 25 % lete~ega pepela (SEM, odbiti elek-
troni). Tanka pu{~ica: delec, bogat s Pb in Sb, debela pu{~ica: spinel
Figure 3: Backscattered electron micrographs of glass samples after
vitrification: a) fly ash without addition of waste glass, b) sample P50
with 50 the mass fraction of % of fly ash and c) sample P25 with 25 %
fly ash. Thin arrow: particle rich in Pb and Sb, wide arrow: spinel
Slika 2: Difraktogrami za lete~i pepel (spodnja krivulja, NaCl, KCl,
-SiO2, CaSO4, CaO, Ca2SiO3Cl2, Fe2O3…) za vzorec P50 s 50 %
lete~ega pepela po zasteklitvi (srednja krivulja, amorfna zgradba) in za
vzorec P50 s 50 % lete~ega pepela po devitrifikaciji (zgornja krivulja,
prevladuje volastonit CaSiO3)
Figure 2: XRD traces for fly ash (bottom curve, NaCl, KCl, -SiO2,
CaSO4, CaO, Fe2O3...); for a vitrified sample P50 with 50 % fly ash
(middle curve, amorphous structure) and for a devitrified sample P50
with 50 % fly ash (upper curve, mainly wollastonite CaSiO3).
V vzorcu, ki je bil zasteklen brez dodatka stekla 20,
sta bila prekora~ena samo parametra vrednosti izlu`ka za
elementa antimon (Sb) in selen (Se), ki verjetno izvirata
iz uporabe narezanih avtomobilskih gum kot alternativ-
nega goriva v se`igalnici.
3.3 Analiza steklokeramike
Na segrevalni krivulji DSC-termograma stekla P50 se
pojavita dva ve~ja eksotermna vrhova (slika 4): prvi med
585 °C in 657 °C ter drugi med 815 °C in 937 °C. To
ka`e, da se devitrifikacija za~ne pri temperaturah okoli
600 °C in da pote~e v veliki meri do temperature 950 °C.
Na osnovi rezultatov smo tako dolo~ili temperaturi prve
in druge stopnje devitrifikacije: za prvo stopnjo 700 °C,
pri kateri naj bi nastalo ve~je {tevilo kristalnih jeder, ter
za drugo stopnjo 950 °C, pri kateri dose`emo pospe{eno
rast kristalov, da devitrifikacija pote~e v primerno krat-
kem ~asu.
Mikrostruktura vzorca s 25 % lete~ega pepela ne
ka`e bistvenih sprememb glede na za~etno steklasto
stanje (slika 5a), vendar pa se lastnosti `e bistveno spre-
menijo, saj se nekoliko pove~a trdota, tla~na trdnost pa
se pove~a kar za dvakrat (tabela 2). To lahko pripi{emo
tvorbi drobnih kristalov v amorfni osnovi, kakor tudi
zmanj{anju notranjih napetosti. Slednje je verjetno
pripomoglo k zmanj{anju raztrosa rezultatov. Po drugi
strani so mikrostrukturne spremembe v vzorcih s 50 %
lete~ega pepela (P50) dobro vidne `e pri manj{ih pove-
~avah (slika 5b). Velika grbina (XRD-analiza, zgornja
krivulja (slika 2)), ki se skoraj v celoti sklada s {irokim
vrhom stekla, ka`e, da vzorec {e ni v celoti kristaliziran.
Uklonski vrhovi se dajo v celoti pojasniti z navzo~nostjo
volastonita CaO × SiO2, ~eprav so v mikrostrukturi {e
drobni delci drugih faz, pa tudi EDS-analiza poka`e
navzo~nost drugih elementov v tej kristalni fazi
Al0,00Ca0,14Cr0,04K0,00Mg0,02Na0,005Si0,20O0,58, tako da
sestava bolj ustreza formuli (Ca,Cr,Mg)O × SiO2,
amorfni del pa se predvsem obogati z elementi, ki se ne
izlo~ajo v kristalni fazi (Na, K, Al).
Rezultati izlu`evanja vzor~ne steklokeramike so
bistveno pod predpisanimi mejnimi vrednostmi za odla-
ganje na odlagali{~ih za stabilizirane in nereaktivne
nevarne odpadke (tabela 2). Tak{ne lastnosti steklokera-
mike omogo~ajo tudi predelavo tega odpadka v koristne
proizvode.
Izmerjena gostote stekla in steklokeramike so bile
okrog 3 000 kg/m3. To je kar 3,6-krat ve~ v primerjavi z
nasipno gostoto lete~ega pepela ((825 ± 7) g/dm3), kar je
zelo pomemben podatek pri odlaganju materialov. ^e
upo{tevamo, da bi v 1 m3 lahko odlo`ili le 825 kg lete-
~ega pepela, bi lahko v isti odlagalni prostor (1 m3)
odlo`ili 3 000 kg taljenega vzorca. Vendar je treba upo-
{tevati, da kon~ni produkt steklo verjetno ne bi bil v
enem kosu, temve~ bi bil me{anica manj{ih ko{~kov
velikosti oreha in bi bila zato njegova nasipna gostota
manj{a. S postopkom taljenja lahko torej izredno zmanj-
{amo prostornino odpadkov, kar omogo~a dolgotrajnej{o
uporabo razpolo`ljivega odlagalnega prostora. Pomemb-
no je tudi dejstvo, da je tla~na trdnost predvsem steklo-
keramike dovolj velika, da se zmanj{a mo`nost drob-
ljenja, s tem pa se zaradi manj{e aktivne povr{ine
zmanj{a tudi stopnja izlu`evanja.
4 SKLEPI
Na osnovi rezultatov raziskav smo pri{li do nasled-
njih sklepov:
– Lete~i pepel in odpadno steklo se v teko~em stanju
dobro me{ata med seboj, zato nastane po litju v
kovinsko kokilo produkt, ki je prete`no amorfen. Del
lete~ega pepela pri taljenju sicer izpari, zato zahteva
njegovo taljenje ustrezne ukrepe (delo v digestoriju),
vendar se kljub temu v zasteklenem produktu zadr`i
ve~ji dele` lete~ega pepela. Pri tem se je izkazalo
bolj ugodno taljenje pri ni`ji temperaturi (1290 °C).
Poleg tega je ugotovljeno, da se ve~ izparljivih
elementov zadr`i v produktu, ~e je dele` lete~ega
pepela manj{i. Dele` kristalne snovi v zasteklenem
produktu se zmanj{uje z ve~anjem dele`a odpadnega
N. ORE[EK ET AL.: IMOBILIZACIJA LETE^EGA PEPELA Z ZASTEKLITVIJO
Materiali in tehnologije / Materials and technology 44 (2010) 6, 373–378 377
Slika 5: Elektronski mikroposnetki steklokeramike a) devitrificiran
vzorec lete~ega pepela z dodatkom 25 % odpadnega stekla (P25), b)
devitrificiran vzorec lete~ega pepela z dodatkom 50 % odpadnega
stekla (P50) (1 – amorfna osnova, 2 – kristalna faza, pu{~ice: delci
bogati s Pb in Sb).
Figure 5: Backscattered electron micrographs of glass-ceramics: a)
devitrified sample with 25 wt.% fly ash (P25), b) devitrified sample
with 50 % fly ash (P50) (1 – amorphous matrix, 2 – crystalline phase,
arrows – Sb- and Pb-rich particles)
stekla. Te`ki elementi niso enakomerno porazdeljeni
v trdni osnovi, temve~ se prete`no nahajajo v obliki
dispergiranih delcev mikrometrske velikosti.
– Devitrifikacija vzorca s 25 % lete~ega pepela je
potekala bistveno po~asneje kot vzorca s 50 % lete-
~ega pepela. Pri devitrifikaciji se je iz amorfne osno-
ve izlo~al prete`no volastonit s pribli`no formulo
(Ca, Cr, Mg)O × SiO2. Pri devitrifikaciji se trdota ni
bistveno spremenila, medtem ko se je tla~na trdnost
pove~ala vsaj za dvakrat, kar lahko delno pripi{emo
tudi relaksaciji notranjih napetosti.
– Z zasteklitvijo lete~ega pepela smo dobili iz nevar-
nega odpadka produkt, ki ga lahko nenevarno odla-
gamo na odlagali{~a. Ne samo to, njegove lastnosti
bi lahko omogo~ile predelavo v uporabne izdelke,
predvsem kot nasipni material. Predelava lete~ega
pepela v Sloveniji za zdaj {e ni ekonomsko upravi-
~ena, vendar se je zaradi pri~akovane podra`itve
odlaganja v prihodnosti in pove~anja koli~in lete~ega
pepela smiselno {e naprej ukvarjati tako z optimi-
zacijo postopka zasteklitve, kot tudi z razvijanjem
novih inovativnih procesov imobilizacije lete~ega
pepela.
ZAHVALA
Zahvaljujemo se In{titutu za tehnologijo materialov
na Fakulteti za strojni{tvo Univerze v Mariboru (pred-
vsem mag. Tonici Bon~ina), Katedri za livarstvo Nara-
voslovnotehni{ke fakultete Univerze v Ljubljani (dr.
Jo`etu Medvedu in dr. Primo`u Mrvarju), Laboratoriju
za strojne elemente in konstrukcije (dr. Nenadu Gube-
ljaku in dr. Jo`etu Predanu) ter vsem, ki so nam poma-
gali pri izlu`evanju (Laboratorij za barvanje, barvno
metriko in ekologijo plemenitenja Fakultete za stroj-
ni{tvo UM Univerze v Mariboru, ERICO Velenje in
Zavod za zdravstveno varstvo Celje (ZZV Celje)).
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N. ORE[EK ET AL.: IMOBILIZACIJA LETE^EGA PEPELA Z ZASTEKLITVIJO
378 Materiali in tehnologije / Materials and technology 44 (2010) 6, 373–378
M. SUBAN ET AL.: THE IMPACT OF STAGNANT WATER ON THE CORROSION PROCESSES ...
THE IMPACT OF STAGNANT WATER ON THE
CORROSION PROCESSES IN A PIPELINE
VPLIV ZASTAJAJO^IH VODA V CEVOVODIH NA KOROZIJSKE
PROCESE
Marjan Suban, Robert Cvelbar, Borut Bundara
Institute of metal constructions, Mencingerjeva 7, 1001 Ljubljana, Slovenia
marjan.suban@imk.si
Prejem rokopisa – received: 2010-05-19; sprejem za objavo – accepted for publication: 2010-08-26
The occurrence of stagnant water is often the result of the incomplete or improper hydrotesting procedure of pipelines (e.g., the
water installation in buildings) or improper design and construction of a pipeline with blind ends. The influence of stagnant
water can also be observed in sprinkler fire-protection systems. Potable water, which is not additionally chemically treated, is
commonly used for the hydrotest or as a medium in fire-protection systems. Water, therefore, contains bacteria that cause
microbiologically influenced corrosion. This relatively unusual form of corrosion results from the interactions of bacteria with
various metals and their alloys and can increase the corrosion rate up to 100-times above that in conventional types of corrosion.
The article describes in detail the causes for the formation and progress of microbiologically influenced corrosion and its
consequences for the case of galvanized water pipes. Some recommendations for a reduction of the risk concerning the
microbiologically influenced corrosion in water pipes are stated in the conclusion.
Keywords: microbiologically influenced corrosion, water pipes, hydrotest
Pojav zastajajo~ih ali mirujo~ih voda je pogosto posledica neupo{tevanje celotnega postopka izvedbe tla~nega preizkusa
cevovodov (npr. vodovodne instalacije v zgradbah) ali nepravilnega projektiranja in izvedbe cevovoda, kjer se pojavljajo slepi
odcepi. Zastajajo~i vpliv vode je mogo~e najti tudi v sprinklerskih sistemih protipo`arne za{~ite. Za izvedbo tla~nega preizkusa
ali kot medij v protipo`arnih sistemih se pogosto uporablja pitna voda, ki ni dodatno kemijsko obdelana in zato vsebuje tudi
bakterije, ki povzro~ajo mikrobiolo{ko vplivano korozijo cevovoda. Ta dokaj nenavadna oblika korozije je rezultat interakcije
med bakterijami in raznimi kovinami ter njihovimi zlitinami, zanjo pa je zna~ilna tudi do 100-krat ve~ja hitrost kot pri navadnih
tipih korozije. V ~lanku so podrobneje opisani vzroki za nastanek ter potek mikrobiolo{ko vplivane korozije ter njene posledice
pri vro~e pocinkanih vodovodnih ceveh. V sklepu je podanih tudi nekaj priporo~il za zmanj{evanje tveganja za pojav
mikrobiolo{ko vplivane korozije v vodovodnih ceveh.
Klju~ne besede: mikrobiolo{ko vplivana korozija, vodovod, tla~ni preizkus
1 INTRODUCTION
In a pipeline, which is basically designed to transport
liquid, gaseous or solid materials (strewn materials),
stagnant or standing water can often occur. Stagnant
water and water with a low flow rate, form a potential
breeding ground for the development of microorganisms.
These not only represent a risk to living organisms, but
they also act on non-living materials, such as various
metals and other materials. Bacteria, fungi and algae are
microorganisms, which among other things can cause an
increase in the corrosion of metals and their alloys. For
this type of corrosion the name Microbiologically
Influenced Corrosion, or MIC for short, is used. Similar
considerations also apply to systems that are shown in
Table 1 and are vulnerable in terms of MIC. The table
Materiali in tehnologije / Materials and technology 44 (2010) 6, 379–383 379
Table 1: Systems with persistent MIC problems 1
Tabela 1: Sistemi, kjer se pojavlja mikrobiolo{ka vplivana korozija 1
Application/System Problem Components/Areas Microorganisms
Pipelines/storage tanks (water,
wastewater, gas, oil)
Stagnant areas in the interior Exterior of
buried pipelines and tanks, especially in
wet clay environments
Inadequate drying after hydrotesting
Aerobic and anaerobic acid producers
SRB
Iron/manganese oxidizing bacteria
Sulfate oxidizing bacteria
Cooling systems
Cooling towers
Heat exchangers
Storage tanks
Aerobic and anaerobic bacteria
Metal oxidizing bacteria
Slime forming bacteria
Algae
Fungi
Vehicle fuel tanks Stagnant areas Fungi
Power generation plants
Heat exchangers
Condensers
Aerobic and anaerobic bacteria
SRB
Metal oxidizing bacteria
Fire sprinkler systems Stagnant areas Anaerobic bacteriaSRB
UDK 624.07:620.1/.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)379(2010)
also presents the microorganisms that are the most
commonly found as a cause of corrosion.
Sulfate-reducing bacteria or SRB are found in a
variety of natural environments. They are the best known
bacteria that have been found to be involved in MIC
problems. The growth of the bacteria occurs in the
environment shown in Figure 1 (left). The bacteria in
stagnant or slow-moving water form a biofilm on the
surface of the metal, under which the metal corrosion
processes run up to 100 times faster than with normal
corrosion 2.
For the occurrence of MIC chemical reactions,
microorganisms are required and fluids with micro-
organisms represent a new source of cathode reactants.
Although there is no universal agreement about the
mechanism by which SRB cause or accelerate metallic
corrosion, it is believed that the bacteria bring about a
cathodic depolarization of the metal surface by removing
hydrogen from the cathodic sites in a sulfate-reducing
reaction. The reactions for this proposed mechanism are
as follows:
Anodic reaction: 4Fe  4Fe2+ + 8e– (1)
Electrolytic dissociation of water:
8H2O  8H
+ + 8OH– (2)
Cathodic reaction: 8H+ + 8e–  8H (3)
Cathodic depolarization by SRB:
8H+ + SO4
2–
 4H2O + S
2– (4)
Corrosion product:
Fe2+ + S2–  FeS (5)
Corrosion product:
3Fe2+ + 6OH–  3Fe(OH)2 (6)
Overall reaction:
4Fe + 4H2O + SO4
2–
 3Fe(OH)2 + FeS + 2OH
– (7)
Therefore, in the case of ferrous alloys SRB cause the
formation of iron sulfide (FeS). In the case of a reaction
with zinc (Zn), which represents the corrosion protection
of steel, the corrosion products are zinc sulfide (ZnS)
and its aggregates, which can be found in biofilm, as is
shown in Figure 1 (right).
In addition to SRB, other types of bacteria that cause
the MIC of metals are known, such as Desulfovibrio,
Gallionella or Pseudomonas. They attack various types
of metals, like iron, stainless steel, aluminum, zinc and
copper 4. Pipelines with a pH of the water between 4 and
8 and a temperature from 20 °C to 45 °C are, in terms of
microbiology, ideal for the growth of microorganisms 5.
In addition to the temperature, the pH and the oxygen
content, the development of bacteria is also affected by
other physico-mechanical factors, such as the surface
energy, the surface roughness of the metal, and the fluid
flow rate. A useful diagram for assessing the risk of MIC
in a pipeline was given by Krooneman et al. 6 (see Figu-
re 2). What is emphasized in this figure is that micro-
organisms, like human beings, are able to live only in
certain conditions, of which the fluid-flow velocity is
extremely important. As a threshold the flow velocity in
Figure 2 is given as 1.5 m/s. Below this velocity the
formation of a bacterial biofilm is possible, but above
this value there are fewer options because the fluid flow
velocity is too high. Johnsen and Bardal mention that
microorganisms may also form a biofilm in a fluid-flow
velocity of about 4.5 m/s 7. Without regard to this
suggestion, it is a fact that by reducing the flow velocity,
the mechanical force that can prevent the formation of
biofilm on the wall of the pipeline is also reduced. The
development of MIC is therefore more likely.
2 EXPERIMENTAL WORK
For this study of the impact of stagnant water on MIC
a real object was used. A water pipeline in a hospital
building, which was completed in 2002, was analyzed.
M. SUBAN ET AL.: THE IMPACT OF STAGNANT WATER ON THE CORROSION PROCESSES ...
380 Materiali in tehnologije / Materials and technology 44 (2010) 6, 379–383
Figure 1: Components of the environment for the development of
MIC (right) and microscopic photograph of biofilm in the case of zinc
corrosion (left) 3
Slika 1: Sestavni deli okolja za razvoj mikrobiolo{ke vplivana koro-
zije (desno) in mikroskopski posnetek biofilma v primeru korozije
cinka (levo) 3
Figure 2: Schematic relationship between the flow velocity and the
pH of water and the temperature T with respect to the corrosion rate R
Slika 2: Shematski prikaz odvisnosti hitrosti korozije R od hitrosti
pretoka fluida v, temperature T in kislosti fluida pH
Figure 3: Corrosion products on the lower part of the pipe in its inte-
rior (left) and the longitudinally cut pipe (right)
Slika 3: Korozijski produkti na spodnje delu cevi v njeni notranjosti
levo ter vzdol`no razrezana cev
Hydrotesting was performed before the acceptance of the
pipeline. For the implementation of the hydrotest ordi-
nary drinking water was used. This water was not addi-
tionally chemically treated. The normal use of the water
pipeline in the object started in 2006. In the meantime,
the system was wholly or partly emptied. Since then,
until the end of 2008, eight heat shocks with a water
temperature higher than 70 °C and several chlorine
shocks were carried out in the water-supply system.
The water supply system in the object was built from
galvanized steel pipes of different diameters with the
chemical composition in mass fractions of the base steel:
C-0.10; Si-0.13; Mn-0.48; P-0.04; S-0.02 and Fe-Ba-
lance. The measured thickness of the galvanized layer of
zinc (hot-dip galvanizing) ranged from 35 μm to 90 μm.
Inspection of the corrosion inside the water pipe was
carried out using a Everest PLS 500 DA videoscope.
Selected parts of the pipes were cut from the system and
further analyzed by macro sections. The corrosion
products were also chemically analyzed using energy-
dispersive X-ray spectroscopy (scanning electron micro-
scope JEOL 5500 LV with an EDX analyzer).
3 RESULTS AND DISCUSSION
3.1 Appearances of corroded sites
When inspecting the pipe interior the first finding
was that it was coated with a carbohydrate biofilm.
Corrosion, which is much more intense in the lower part
of the pipe (see Figure 3), shows that the corrosion
process started mainly due to stagnant water and an
incompletely drained water-supply system. This shows
the explicit effect of stagnant water, probably after the
hydrotest, or at another time, when the system was, for a
longer period, only partially drained. The smaller quan-
tity of stagnant water was a very appropriate medium for
the growth of microorganisms and the further develop-
ment of MIC. The microorganisms dissolve zinc and iron
and even lower-quality stainless steel (e.g., 18Cr-8Ni
austenitic stainless steel AISI 304) does not passivate in
such a medium 8.
The color of the corrosion products is an essential
indicator in the investigation of MIC. The reddish-brown
deposits with a black corroded surface in Figure 4 may
be good indicators of the SRB and iron-oxidising bac-
teria. The black corroded surface under the carbohydrate
biofilm was an indication of the iron sulfide formation. A
chemical analysis of the corrosion products was not
performed. A distinctive smell of rotten eggs, which
indicates hydrogen sulfide gas produced by the SRB,
arises during the removal of the biofilm.
A classification of the morphological types of surface
corroded sites was made, but on this basis it was not
possible to make a single conclusion. Similar data have
been obtained by other authors, e.g., in 9, where it was
found that the corrosion surface morphology was related
to the chemistry at the metal surface and not to the
presence or absence of microorganisms.
At the next stage the study was extended with a
metallographic examination of the cross sections of
corroded sites. The researchers in 10 described ink-bottle-
shaped pits in the stainless steel as diagnostic of MIC
(Figure 5 left). This, in our case (Figure 5 right), was
not confirmed, but it is true that in our case the material
was galvanized carbon steel. After the breakthrough of
the zinc protective layer, the corrosion of the underlying
steel started. It is possible that the further growth in the
depth of the corrosion in the carbon steel could develop a
similar form as in Figure 5 (left).
3.2 EDX analysis
The corrosion products formed in the pit were
analyzed with the EDX method and the graph in Figure
6 shows peaks for iron (Fe), zinc (Zn), oxygen (O) and
sulfur (S). As already shown by Equation 7, in the MIC
iron sulfide (FeS) was formed as a corrosion product. For
this reason, the EDX analysis of the corrosion products
detected some amount of sulfur (w = 0.42 %), which is
characteristic of this type of corrosion. The presence of
sulfur in conjunction with the zinc was also suggested by
the presence of zinc sulfide (ZnS) in the corrosion
products. In addition, we have also found a significant
amount of chlorine in the corrosion products, probably in
the form of Cl- ions. Thus, the large amount of chlorides
is probably due to the implementation of chlorine shocks
in the water supply system. The effect of chlorine on the
creation of new corrosion pits and the deepening of the
M. SUBAN ET AL.: THE IMPACT OF STAGNANT WATER ON THE CORROSION PROCESSES ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 379–383 381
Slika 4: Korozijski produkti na povr{ini cevi
Figure 4: Corrosion products on the surface of the pipe
Slika 5: Pre~ni presek korozijske po{kodbe levo tipi~na pri nerjavnih
jeklih 10, desno v na{em primeru pri pocinkanju na notranji strani cevi
Figure 5: Cross-section illustration of an ink-bottle-type pit noted in
the case of stainless steel (left) 10 and the cross-section of the pit in the
investigated case of galvanized carbon steel (right)
already formed pits due to MIC, was not the subject of
this investigation.
For a comparison, we also analyzed a water pipe on
which we could only find traces of white corrosion
(Figure 7, top right). From the analysis of the corrosion
products on the galvanized surface we found that the
corrosion products consisted only of the products of zinc
with the oxygen bound as zinc hydroxide (Zn(OH)2)
(Figure 7). We have also found traces of iron. As the
proportion of iron in these corrosion products was very
small, we concluded that it originated from the lower
layers of zinc coating in which there can be from 7 to
about 20 weight percent of iron (Fe) in the form of
various intermetallic phases.
3.3 Water-quality analysis
In this investigation we also analyzed the water
quality of the water-supply system, which in operation
was subject to MIC. Ten samples of water were taken
and analyzed using flame atomic absorption spectro-
metry (FAAS) and inductively coupled plasma mass
spectroscopy (ICP-MS) to determine the concentration
of the dissolved iron and zinc. Two samples were taken
at the terminal to the city water-supply system and show
that iron and zinc were not from the city water-supply
system. Among eight samples of water originating from
the observed water-supply system, an increased content
of iron up to 38 mg/L was found. Similarly, the content
of zinc was above 26 mg/L. The results are comparable
to a similar analysis of water under laboratory conditions
in 11, where the values of the zinc content in the water
amounted to about 28 mg/L in samples with MIC. The
measured proportion of the zinc and iron contents in the
water are higher than the permissible and/or
recommended limits for drinking water, which are 0.2
mg/l for iron and 3 mg/l for zinc 12.
4 CONCLUSIONS
The present paper describes the occurrence of MIC in
galvanized steel on a real object, as a result of the
activity of microorganisms in stagnant water. The results
of the analyses of the corrosion products, especially the
EDX analyses, show that in the corrosion products there
are both elements and compounds (e.g., sulfur in the
form of FeS and ZnS) reflecting the activity of the
microorganisms. The corrosion products found originate
from corrosion due to SRB and iron-oxidizing bacteria,
the presence of which was not proved with a biological
analysis.
It is concluded that it is necessary to avoid the
occurrence of stagnant water in water-supply systems,
because it represents a potential source of micro-
organisms leading to the development of MIC. The
emergence of stagnant water is often the result of irregu-
larities in the construction, the improper execution of the
hydrotest or because of the system itself (e.g., fire-
protection sprinkler systems). To avoid the occurrence of
MIC, some recommendations need to be considered:
• When designing and constructing a water-supply
system a blind branch should be avoided. The design
of horizontal pipelines needs to be "self-draining".
• The pipeline must be constructed so that the velocity
of the fluid flow is at least 1.5 m/s.
• Only chemically treated water should be used to fill
the fire-protection sprinkler system.
• Demineralized drinking water should be used to carry
out the hydrotest. As soon as possible after the com-
pletion of the hydrotest, it is necessary to drain and
dry the pipeline.
5 REFERENCES
1 P. J. Scott; S. Borenstein, F. Blackburn; B. Cookingham; D. De-
marco, Expert consensus on MIC: Prevention and monitoring, Part 1,
Materials Performance, 43 (2004) 3, 50–54
2 G. J. Licina, G. Nekoksa, An innovative method for rapid detection
of microbiologically influenced corrosion, Tri-Service Corrosion
Conference, 1994, Orlando FL, 217–229
3 M. Labrenz, G. K. Druschel, T. Thomsen-Ebert, B. Gilbert, S. A.
Welch, K. M. Kemner, G. A. Logan, R. E. Summons, G. De Stasio,
P. L. Bond,. B. Lai, S. D. Kelly, J. F. Banfield, Formation of
sphalerite (ZnS) deposits in natural biofilms of sulfate-reducing
bacteria, Science, 290 (2000) 5497, 1744–1747
M. SUBAN ET AL.: THE IMPACT OF STAGNANT WATER ON THE CORROSION PROCESSES ...
382 Materiali in tehnologije / Materials and technology 44 (2010) 6, 379–383
Slika 6: EDX analiza produktov mikrobiolo{ke vplivane korozije na
povr{ini pocinkanega jekla
Figure 6: EDX analysis of the corrosion products formed on the
galvanized steel surfaces with microbiologically influenced corrosion
Slika 7: EDX analiza produktov in posnetek dela cevi (zgoraj desno) z
belo korozijo pocinkanega jekla
Figure 7: EDX analysis of the corrosion products and photograph of
pipe section (top right) with the white corrosion of galvanized steel
4 S. C. Dexter, Microbiologically influenced corrosion, corrosion: fun-
damentals, testing, and protection, ASM Handbook, vol 13A, ASM
International, (2003), 398–416
5 R. Javaherdashti, Microbiologically influenced corrosion: An Engi-
neering Insight, Springer London, 2008
6 J. Krooneman, P. Appeldoorn, R. Tropert, Detection prevention and
control of microbial corrosion, Proceedings of Eurocorr Maastricht,
2006
7 R. Johnsen, E. Bardal, Cathodic properties of different stainless
steels in natural seawater, Corrosion, 41 (1985) 5, 296–302
8 I. Kos, F. Gre{ovnik, M. Gre{ovnik, Korozijska obstojnost nekaterih
nerjavnih jekel = The corrosion resistance of stainless steel, Mater.
Tehnol., 36 (2002) 5, 271–273
9 R. E. Tatnall, D. H. Pope, Identification of MIC. Ch. 8, In: A prac-
tical manual on microbiologically influenced corrosion, Kobrin G
(ed), NACE, Houston, Texas USA, 1993
10 B. J. Little, J. S. Lee, R. I. Ray, Diagnosing microbiologically influ-
enced corrosion: A state of the art review, Corrosion 62 (2006) 11,
1006–1017
11 E. Ilhan-Sungur, N. Cansever, A. Cotuk, Microbial corrosion of gal-
vanized steel by a freshwater strain of sulphate reducing bacteria
(Desulfovibrio sp.), Corrosion Science 49 (2007) 3, 1097–1109
12 World Health Organization, Guidelines for Drinking-water Quality,
Third edition incorporating the first and second addenda, Vollmer,
2008, 390 and 459
M. SUBAN ET AL.: THE IMPACT OF STAGNANT WATER ON THE CORROSION PROCESSES ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 379–383 383

D. KOZAK et al.: DETERMINATION OF THE CRITICAL PRESSURE ...
DETERMINATION OF THE CRITICAL PRESSURE FOR A
HOT-WATER PIPE WITH A CORROSION DEFECT
DOLO^ITEV KRITI^NEGA PRITISKA V VRO^EVODNI CEVI S
KOROZIJSKO PO[KODBO
Dra`an Kozak, @eljko Ivandi}, Pejo Kontaji}
J. J. Strossmayer University of Osijek, Mechanical Engineering Faculty in Slavonski Brod, Trg Ivane Brli}-Ma`urani} 2,
HR-35000 Slavonski Brod, Croatia
drazan.kozak@sfsb.hr
Prejem rokopisa – received: 2009-10-21; sprejem za objavo – accepted for publication: 2010-08-31
This paper deals with a determination of the critical pressure, i.e., the allowable operating pressure, of a corroded pipeline. This
allowable pressure, as well as the safe-working pressure, are determined by using the analytic expressions given in the
DNV-RP-F101 procedure for corroded pipelines. The failure pressure was calculated numerically using the finite-element
method (FEM) with assumption of linear-elastic and ideal-plastic behaviour of the pipe material.
Key words: corroded hot-water pipe, defect, DNV-RP-F101 procedure, allowable pressure, FEM analysis
V delu obravnavamo dolo~itev kriti~nega dovoljenega delovnega pritiska v korodirani cevi. Dovoljeni pritisk in varen
obratovalni pritisk sta bila dolo~ena z uporabo analiti~nih izrazov v DNV-RP-F101-proceduri za korodirane cevovode. Raztr`ni
pritisk je bil izra~unan numeri~no po metodi kon~nih elementov (FEM) s predpostavko o linearnem elasti~nem in idealnem
plasti~nem vedenju materiala cevi.
Klju~ne besede: korodirana vro~evodna cev, po{kodba, DNV-RP-F101-procedura, dovoljeni pritisk, FEM-analiza
1 INTRODUCTION
Cylindrical shells under pressure are extensively used
in various industrial structures including pipelines for
oil, gas and hot-water transport. Since such usage usually
requires underground exploitation for a longer period of
time, such pipelines are subjected to damage 1 under
external environmental conditions (primarily corrosion)
as well as under mechanical factors (including loadings
of structure). Such damage can lead to initiation and
growth of the surface crack and finally lead to failure of
structure 2. Therefor, a need for investigation of influence
of structural damage of cylindrical shells is present.
Various studies have been done dealing with this
problematic, including recent ones 3-6, using analytical
and numerical approach to investigate effect of wall
thinning, crack initiation and growth in pipes under
internal pressure. Because of the importance of described
problematic, this paper also deals with investigation of
corrosion effects on pressurized cylindrical shell which
is part of one pipeline of the hot water system.
2 CORROSION DEFECT OF HOT-WATER PIPES
The corrosion defects in hot-water pipes of the city
system for hot-water distribution in Osijek, Croatia were
analysed. Since the material of pipes in system is usually
a high-quality unalloyed steel, which is not suitable for
heat treatment, it can be affected by corrosion during
exploitation. The corrosion mass-loss decreases the
cross-section of the pipe at the point of damage. To
quantify the effect of wall thinning due to high corrosion
effect and to avoid eventual pipe failure (leaking of the
pipe shown on Figure 1), a maximum allowable opera-
ting pressure of the pipe had to be determined.
The pressures in the hot-water pipes of the hot-water
system in the city of Osijek are from 0.8 bar to 13 bar (in
normal exploitation), and in the temperature range is
70–140 °C. All the fittings and pipes are made for
normal pressure NP 16 (16 bar).
In this investigation, a control calculation for the
pipeline (with NO 250 opening) with wall thickness 5
mm and with the measured geometry of corrosion
Materiali in tehnologije / Materials and technology 44 (2010) 6, 385–390 385
UDK 621.791:620.193:62-462 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)385(2010)
Figure 1: Leaking of hot-water pipes during use
Slika 1: Pu{~anje vro~e vode med obratovanjem
defects found during exploitation of the hot water system
of Osijek is done. The Det Norske Veritas DNV-RP-F101
procedure (for corroded pipes 7) was used for making
analytic calculations and the defects were modelled
numerically with the finite element method (FEM) 8. The
values of allowable pressure according to both proce-
dures were determined and compared for different defect
geometries.
Several kilometres of pipelines (of different dimen-
sions from NO40 to NO350) were repaired this year and
it was possible to determine and measure in-situ the real
damage of pipes. The pipe with a characteristic defect, as
well as the detail of defect size determination, is shown
in Figure 2.
3 ANALYTICAL DETERMINATION OF THE
ALLOWABLE PRESSURE OF A CORRODED
PIPELINE ACCORDING TO THE DNV-RP-F101
PROCEDURE
Det Norske Veritas recommended the procedure for
calculating the critical value of the pipeline pressure with
determined defects in the document DNV-RP-F101 7,9.
There are also some limitations during the application:
– carbon steel has to be used for making pipes,
– cyclic loadings and defects (e.g., cracks) are not
considered,
– corrosion and mechanical defects are not combined,
– it is presumed that there are no defects in the pipe
welds,
– it can be applied only up to the depth of the defect,
and not greater than 85 % of the wall thickness.
It is also assumed that modern steels for pipes have
an adequate toughness and it can be expected that the
so-called plastic collapse of the pipes will occur. This
collapse occurs when the equivalent stresses exceed the
yield point through the whole remaining ligament in
front of the defect (when looking through the thickness
of the pipe).
The analysis in this paper is limited to the assumption
of a single defect and interacting defects are not
considered. The calculation was performed for the pre-
sence of an outer, longitudinal corrosion defect and the
pipe was loaded with an internal pressure. The super-
imposing of the effects of more close defects, the axial
forces and the bending or temperature were not
considered. The influence of an error in the radial
direction was also not considered. It is assumed that the
dimension of the defect in the radial direction is smaller
than in the axial one. For calculation purposes, the
irregular shape of the corrosion defect is idealised to
rectangular (Figure 3).
Two approaches are possible for assessing the
integrity of the corroded pipes and the main difference is
based on a different safety philosophy. The first
approach is probabilistic and it includes the safety
factors that consider possible deviations of mechanical
properties of the material and changes in wall thickness,
i.e., the internal pressure. In this way, the measurement
uncertainty of the defect dimensions and the material
property specification are included in the determination
of the allowable operating pressure. The second
approach is based on the ASD (allowable stress design)
format. The allowable pressure (capacity) of the pipeline
with the corrosion defect is calculated, and this failure
pressure is divided by a safety factor based on the
original design factor. When assessing the corrosion
defects, due consideration should be given to the
measurement uncertainty of the defect dimensions and
the pipeline geometry. These uncertainties are not
included in the second approach, and are left to the user
to consider and account for in the assessment.
3.1 Determination of the allowable operating pressure
of a corroded pipeline
The maximum allowable operating pressure for a
pipeline with a corrosion defect (metal loss) with an
internal pressure is given by the acceptance equation:
p
t f
D t
d
t
d
t
u
corr m
d
d
= ⋅
⋅ ⋅
−
⋅
− ⋅ ⎛⎝
⎜ ⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟
−
⋅ ⎛⎝



2
1
1
( )
*
⎜ ⎞
⎠
⎟
⎛
⎝
⎜
⎜
⎜⎜
⎞
⎠
⎟
⎟
⎟⎟
*
Q
(1)
where:
D. KOZAK et al.: DETERMINATION OF THE CRITICAL PRESSURE ...
386 Materiali in tehnologije / Materials and technology 44 (2010) 6, 385–390
Figure 2: Determining the geometry of hot-water-pipe defects
Slika 2: Dolo~itev geometrije po{kodb vro~evodne cevi
Figure 3: Illustration of irregular and rectangular defects
Slika 3: Oblika neenakomerne in pravokotne po{kodbe
m Partial safety factor for the longitudinal corrosion-
model prediction,
t Uncorroded, measured, pipe-wall thickness, in mm,
fu Tensile strength to be used in the design, in N/mm2,
D Nominal outside diameter, in mm,
d Partial safety factor for the corrosion depth,
Q Length correction factor, in mm.
It is important to point out that the partial safety
factors m and d, as well as the ratio (d/t)* depend on the
applied method of inspection, i.e., the accuracy of the
measured defect depth. In this case, the safety factor m
was determined for the low class of measuring safety as
m = 0,79. Furthermore, the measured (d/t)* ratio is
defined as:
d
t
d
t
d
t
⎛
⎝
⎜ ⎞
⎠
⎟ = ⎛⎝
⎜ ⎞
⎠
⎟ + ⋅ ⎡⎣⎢
⎤
⎦⎥
*
meas
d StD (2)
where:
(d/t)meas Measured (relative) defect depth,
d Factor for defining a fractile value for the corrosion
depth
StD[d/t] Standard deviation of the measured (d/t) ratio.
In this paper, it is assumed (with the permission of 7),
that d = 2, i.e., StD[d/t] = 0.16 with d = 1.2. The
length-correction factor is defined as:
Q
D t
= + ⋅
⋅
⎛
⎝
⎜ ⎞
⎠
⎟1 031
1
2
. (3)
The pipe material was unalloyed carbon steel St. 37.0
(according to DIN 1629) with less than 0.16 % C, with a
yield point ReH = 235 MPa, i.e., tensile strength Rm =
360–440 MPa and elongation A5 = 25 %. The lower
value of tensile strength, i.e., fu = 360 MPa, was consi-
dered in the calculation.
The outside diameter of the pipe is D = 273 mm and
the pipe-wall thickness is t = 5 mm. The ratio of the
corrosion defect depth d and the pipe-wall thickness t is
as follows: d/t = (0.1, 0.2, 0.3, 0.4 and 0.5). The defect
length l was also changed by the introduction of the
geometry factor k:
k
D t
=
⋅
1
(4)
where k = (0.2, 0.4, 1.0, 1.84 and 2), which corresponds
to a defect length l = (7.4, 14.8, 36.9, 68 and 73.9) mm.
The allowable pressure, which considers the measure-
ment uncertainty, was calculated and the results are
given in Figure 4.
The allowable operating pressure pcorr = 5.66 MPa for
the measured defect geometry on the pipes (t = 1.5 mm
and l = 68 mm) is shown in Figure 4. It is obvious from
these results that the allowable pressure is almost
constant for the smaller defect lengths (l  10 mm) and
the defect depths up to 1/3 of the pipe-wall thickness (d
 1/3 t).
As a comparison, the allowable operating pressure
for the same geometry is determined acording ASME
code 10 for vessels under internal pressure with expres-
sion (5) giving maximum allowable pressure of 6.03 bar.
P
SEt
R t
=
+ 06.
(5)
where:
R Inner radius of pipe, in mm,
S Allowable stress, in MPa,
E Material and pipe construction quality factor.
3.2 Determination of the failure pressure of the
corroded pipeline
The failure pressure of the corroded pipe with a
single defect can be calculated from:
P
t f
D t
d
t
d
t Q
f
u=
⋅ ⋅
−
⋅
−⎛⎝
⎜ ⎞
⎠
⎟
−
⋅
⎛
⎝
⎜
⎞
⎠
⎟
2
1
1
( )
(6)
with the equation elements the same as for equations (1)
and (3). The pressure where the failure of pipeline
occurs are shown as a diagram (Figure 5). By com-
paring the pressure of the leaking for the measured
defect (d = 1.5 mm and l = 68 mm) with the results for
the allowable pressure acquired by the probabilistic
approach, it should be noted that these pressure values
are two times higher.
Since we are dealing with the pressures where the
failure occurs, we must multiply them by certain safety
factors in order to have a safe working pressure. The
value acquired from (6) must be multiplied by the
modelling factor F1 and the operational usage factor F2:
P F F Psw f= ⋅ ⋅1 2 (7)
According to 7, the following safety factors were
taken:
F1 = 0.9
D. KOZAK et al.: DETERMINATION OF THE CRITICAL PRESSURE ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 385–390 387
Figure 4: The allowable operating pressure of the corroded pipeline
Slika 4: Dovoljen delovni pritisk korodiranega cevovoda
F2 = 0.72
so that Psw = 0.9 · 0.72 · Pf = 0.648 · Pf, i.e., approxima-
tely 65 %, i.e., the same as when the failure pressure is
divided by the safety factor of 1.54. Figure 6 shows the
pressure that allows safe operation of the pipeline, even
with corrosion damage present.
The safe working pressure of the pipeline, calculated
in this way, is higher than the one acquired with the
probabilistic approach. Therefore, the calculation using
the finite-element method was also made.
4 NUMERICAL DETERMINATION OF THE
FAILURE PRESSURE OF THE CORRODED
PIPELINE
The measured geometry of the pipe (defect length =
68 mm, defect depth = 1.5 mm) was used for numeric
modelling of the corrosion defect. Corrosion defect are
irregular, therefore to investigate such defects using
finite element method a crack shape idealization is
required. It is important that crack idealization is done in
such a manner that it yields to conservative results e.g.,
idealized crack has to be more dangerous than the real
one. Since longitudinal cracks in cylindrical shells are
more dangerous than circumferential cracks, the corro-
sion defect is idealized as a longitudinal semi-elliptic
crack (Figure 7). Length of semi-elliptic crack is deter-
mined from maximum longitudinal length of corrosion
damage and depth of crack corresponded to measured
depth of damage. Due to the present symmetry, only ¼
of the pipe was modelled with the finite elements
(Figure 8).
The FEA Crack programme 11 was used for the crack
modelling. It allows more control over the finite element
mesh, especially in the crack region. For the finite
element analysis the model is meshed with isoparametric
finite elements (Figure 9). Finite element mesh con-
sisted of approximately 25.000 nodes. The magnified
D. KOZAK et al.: DETERMINATION OF THE CRITICAL PRESSURE ...
388 Materiali in tehnologije / Materials and technology 44 (2010) 6, 385–390
Figure 7: Idealization of irregular corrosion defect with a semi-
elliptic crack
Slika 7: FEM model cevi s semielipti~no razpoko kot korozijsko po-
{kodbo
Figure 6: The safe working pressure of the corroded pipeline
Slika 6: Varen obratovalni pritisk korodiranega cevovoda
Figure 9: FEM model of the pipe with a semi-elliptic crack as the
corrosion defect
Slika 9: FEM model cevi s semielipti~no razpoko kot korozijsko po-
{kodbo
Figure 5: The failure pressure of the corroded pipeline
Slika 5: Raztr`ni pritisk korodiranega cevovoda
Figure 8: 3D model of the 1/4 of the pipe
Slika 8: Pove~an detajl FEM modela
detail of the finite element mesh is shown in Figure 10.
The real stress-strain diagram is idealized, and the
material is defined as linear-elastic and ideal-plastic.
The final element mesh was imported to the
programme for finite element analysis – ANSYS 10.0.
The model is loaded with internal pressure in order to
determine failure pressure, which causes plastic collapse
of the pipe. The pressure was gradually increased to
allow precise critical load determination. The criterion
for pipe failure was the internal pressure, which causes
equivalent stress through the remaining ligament of the
pipe, higher than the yield point of the material. In this
case, this value was 8.2 MPa for the defined geometry of
the pipe and the crack. Figures 11 and 12 show the
distribution of the equivalent stresses at the point of
plastic collapse of the pipe. The red area with equivalent
stress higher than 235 MPa spreads through the remain-
ing pipe-wall thickness in front of the crack.
5 CONCLUSION
The aim of this paper was to determine critical, i.e.,
the allowable, operating pressure of the pipeline (Ø 273
× 5) mm with the measured geometry of the corrosion
defect (for the hot water supply system of Osijek). The
Det Norske Veritas DNV-RP-F101 procedure was used
for analytic calculations of the corroded pipes, while the
defects were modelled numerically using the finite
element method. The values of the allowable pressure,
i.e., the safe working pressure, were determined and
compared according to both procedures for different
defect geometry. The failure pressure, which causes
plastic collapse of the pipe, was also determined using
the FEM. The probabilistic approach is the most con-
servative for the corrosion defects with depth of 1.5 mm
and length of 68 mm, because it gives the smallest
allowable pressure (5.66 MPa), while the safe working
pressure is 6.85 MPa (6.05 MPa according to ASME
code), if the possible measurement uncertainty of the
defect dimensions are not considered. The numerical
calculation models the defects as a semi-elliptic crack
and gives the failure pressure, which is 8.2 MPa. This
failure pressure is of great practical value when assessing
the hot-water system's integrity by using e.g., a FAD
diagram.
Acknowledgements: The authors would like to
thank Mr Tomislav Ba{kari}, a student of Mechanical
Engineering Faculty in Slavonski Brod, University of
Osijek, Croatia, for his help with the finite-element
modelling.
6 REFERENCES
1 Reliability and risks assessments with water, oil and gas pipelines.
In: Pluvinage G, Elwany MH, editors. NATO science for peace and
security series. Springer; 2008. 349
2 H. Moustabchir, Z. Azari, S. Hariri, I. Dmytrakh, Experimental and
numerical study of stress–strain state of pressurised cylindrical shells
with external defects, Engineering Failure Analysis 17 (2010),
506–514
D. KOZAK et al.: DETERMINATION OF THE CRITICAL PRESSURE ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 385–390 389
Figure 12: Enlarged detail of the field of equivalent stresses in front
of the crack
Slika 12: Pove~an detajl polja ekvivalentnih napetosti na ~elu razpoke
Figure 11: Field of equivalent stresses in front of the crack for a
failure pressure of 8.2 MPa
Slika 11: Polje ekvivalentnih napetosti pred razpoko pri raztr`nem
pritisku 8,2 MPa
Figure 10: Magnified detail of the FEM model
Slika 10: Pove~an detail FEM modela
3 M. Kamaya, T. Suzuki, T. Meshii, Normalizing the influence of flaw
length on failure pressure of straight pipe with wall-thinning,
Nuclear Engineering and Design 238 (2008), 8–15
4 M. Chiodo, C. Ruggieri, Failure assessments of corroded pipelines
with axial defects using stress-based criteria: Numerical studies and
verification analyses, International Journal of Pressure Vessels and
Piping 86 (2009), 164–176
5 D. S. Cronin, R. J. Pick, Prediction of the failure pressure for
complex corossion defects, International Journal of Pressure Vessels
and Piping 79 (2002), 279–287
6 B. Skallerud, E. Berg, K. R. Jayadevan, Two-parameter fracture
assessment of surface cracked cylindrical shells during collapse,
Engineering Fracture Mechanics 73 (2006), 264–282
7 Recommended practice DNV-RP-F101, Corroded pipelines, Det
Norske Veritas, October 2004
8 ANSYS Release 10.0, User’s manual, 2005
9 R. Koers, Corrosion Procedure and Case Studies, 2nd FITNET TN
Training Seminar and Workshop, University of Maribor, October
2004
10 ASME. ASME Boiler and pressure vessel code Section VIII,
Division 1, New York, NY: The American Society of Mechanical
Engineers; 2003
11 FEA Crack, 3D Finite element software for cracks, Version 3.1
Note
The responsible translator for the English language is
B. A. @eljka Rosandi} (Lecturer), J. J. Strossmayer Uni-
versity of Osijek, Mechanical Engineering Faculty in
Slavonski Brod, Trg Ivane Brli}-Ma`urani} 2, HR-35000
Slavonski Brod, Croatia.
D. KOZAK et al.: DETERMINATION OF THE CRITICAL PRESSURE ...
390 Materiali in tehnologije / Materials and technology 44 (2010) 6, 385–390
A. F. ACAR ET AL.: EFFECTS OF VARIATIONS IN ALLOY CONTENT AND MACHINING PARAMETERS
EFFECTS OF VARIATIONS IN ALLOY CONTENT AND
MACHINING PARAMETERS ON THE STRENGTH OF
THE INTERMETALLIC BONDING BETWEEN A DIESEL
PISTON AND A RING CARRIER
VPLIV SPREMEMB V SESTAVI ZLITINE IN PARAMETROV
OBDELAVE NA TRDNOST INTERMETALNE VEZAVE MED
DIESELSKIM BATOM IN NOSILCEM OBRO^KA
Ahu Fahriye Acar1, Fahrettin Ozturk2, Mustafa Bayrak2
1Mopisan Piston Liner Ring Factory, Konya, Turkey
2Department of Mechanical Engineering, Nigde University, 51245, Nigde, Turkey
fahrettin@nigde.edu.tr
Prejem rokopisa – received: 2010-04-30; sprejem za objavo – accepted for publication: 2010-08-24
The intermetallic bonding between an aluminium alloy piston and a cast-iron ring carrier must be able to withstand prolonged
exposure to the high-pressure and high-temperature environment of a diesel engine. Weak cohesion may result in the debonding
of the ring carrier. Such debonding causes devastating damage to the piston while the engine is in service. Thus, the quality and
strength of the intermetallic bonding is very important in piston manufacturing.
In the present study, the effects of variations in the alloy content and the machining parameters on the overall strength of the
intermetallic bonding are investigated. The defects are determined by ultrasonic inspection.
The results indicate that the variation in the alloy content and graphite flakes affects the bond quality adversely. It was
determined that the graphite size is a very important parameter. If the graphite size gets smaller, then stronger intermetallic
bonding is observed. The graphite accumulation can be resolved by choosing the right amount of alloying elements. The
machining process parameters were found to be ineffective for the formation of defects on the intermetallic bonding.
Keywords: diesel piston; ring carrier; Alfin bond, intermetallic bonding
Intermetalma zveza med batom iz aluminijeve zlitine in lito`eleznim nosilcem obro~ka mora dolgo prena{ati obremenitev pri
visoki temperaturi dieselskega motorja. [ibka kohezija lahko povzro~i odstopanje nosilca obro~ka in povzro~i veliko po{kodbo
pri obratovanju motorja. Zato sta kakovost in trdnost intermetalne zveze zelo pomembni pri izdelavi batov. V tem delu smo
raziskali vpliv sprememb v sestavi in parametrov obdelave na trdnost intermetalne zveze. Napake smo odkrili z ultrazvo~no
kontrolo. Rezultati ka`ejo, da spremembe v sestavi zlitine in grafitni listi~i vplivajo razli~no na kakovost zveze. Velikost
grafitnih listi~ev je zelo pomemben parameter in pri manj{ih listi~ih grafita je trdnost povezave ve~ja. Kopi~enje grafita se lahko
prepre~i s pravo vsebnostjo legirnih elementov. Parametri obdelave ne vplivajo na nastanek napak intermetalne zveze.
Klju~ne besede: dieselski bat, nosilec obro~ka, zveza Alfin, intermetalna zveza
1 INTRODUCTION
Pistons are commonly made of a cast aluminium
alloy for reasons of its excellent and lightweight thermal
conductivity. An aluminium-silicon alloy is commonly
used for the piston material and provides the best overall
balance of properties 1,2. The piston features include the
piston head, the piston pin bore, the piston pin, the skirt,
the ring grooves, the ring lands, and the piston rings. A
ring groove is a recessed area located around the
perimeter of the piston that is used to retain the piston
ring. Piston rings are used to prevent the leakage of
combustible gas and to control the oil film of the liner
when the piston is working in the cylinder. The rings are
inserted into the ring grooves on the piston. These
grooves are called the ring carrier. The ring carrier is
made of Ni-resist (Niresist) cast iron consisting of
graphite in a matrix of austenite. The cast iron retains the
integrity of its original shape under heat, load, and other
dynamic forces. These materials usually contain alloying
elements such as nickel, chromium, copper or molyb-
denum to increase the strength or in order to facilitate the
formation of austenitic cast iron. The corrosion and wear
resistances of these materials are quite high. These
materials are suitable for corrosive environments, such as
sour well oils, salts, salt water acids, and alkalis. Niresist
is denser than gray or ductile irons and has a higher
coefficient of thermal expansion. Ring carrier materials
are melted in induction furnaces in the temperature range
1400–1450 °C. Then, the final shape of the ring carrier is
produced by a spin casting machine. These ring carriers
are placed in the piston mould and the molten piston
material is poured in. The most important thing is that
the Niresit die cast material (ferrous) must be bonded by
a non-ferrous piston material during the casting of the
piston. Because of the fact that diesel engines are worked
at high temperatures and under high pressure, the
durability of the bond between the aluminium piston and
the Niresist piston ring is very crucial. The serious
problems that can be caused by bond defects are
Materiali in tehnologije / Materials and technology 44 (2010) 6, 391–395 391
UDK 621.436 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)391(2010)
available in the literature 3. The intermetallic bonding
between the ring carrier and piston must be able to
withstand prolonged exposure to the high-pressure and
temperature environment of the diesel engine. Weak
cohesion may result in the debonding of the ring carrier.
This debonding causes devastating damage to the piston
while the engine is in service. For the aforementioned
reasons these processes should be controlled precisely.
So, the surfaces of the ring carrier material are shot
blasted using small balls. This process is the pre-process
of an alfin (AlFin) bath for better bonding quality. The
AlFin bath is a method for preparing a ferrous surface
for intermetallic bonding 4. The alfin bonding process is
commonly used to bond a nonferrous Al alloy and a
ferrous alloy, such as aluminium diesel engine pistons
and Niresist ring carriers, directly together. It is well
known that cast iron contains carbon as a result of the
casting process. The carbon precipitates out of solution
and is present in the piece as graphite flakes on cooling.
The size and the distribution of the graphite flakes can
vary as a function of the alloy. "It is well known that
certain elements, when combined with iron containing
carbon, act as retarders, that is, impede the graphitization
of the carbon, and that others act as accelerators, that is,
promote graphitization, during the malleableizing
process. For example, sulphur, manganese and chromium
are retarders, while aluminium, nickel, and copper, in
small quantities, are accelerators of the phenomena of
graphitization 5." It is also a parameter of the casting and
cooling processes. The alfin bonding process begins with
the growth of an Al alloy intermetallic surface layer on
the ring carrier by immersion of the ring carrier in a
molten Al alloy bath. The Al alloy bath usually includes
Al and 6 % Si. The immersion time is determined
according to the required layer thickness. The Al alloy
intermetallic layer is typically grown to about 50 μm or
more in thickness. Bath temperatures typically range
from 700 °C to 750 °C. The thickness of the inter-
metallic layer increases with the alfin bath temperature
and immersion time. The immersion time ranges from 3
min to 18 min. The alfin bond is a real bond, which has a
chemical composition close to the FeAl3 and is formed
by the Fe and Al alloy.
As mentioned previously a diesel piston is produced
by casting. Firstly, Al alloys are melted in an induction
furnace. The temperatures are 780 °C for a 12 % Si alloy,
800 °C for an 18 % Si alloy, and 850 °C for a 24 % Si
alloy. The modification and gas-removing processes are
applied to the molten Al alloy and then the molten alloy
is taken to the waiting furnace. The molten Al-Si alloy is
poured into the mould over the ring carrier. This ring
carrier is observed to embed with the Al alloy after
solidification. Then the piston is subjected to machining
to create the final shape.
Ultrasonic inspection techniques are widely used in
an industry where a non-destructive evaluation of a
workpiece is required. The ultrasonic imaging technique
is founded on pulse-echo techniques: a transducer emits
a short pulse of ultrasonic waves towards the sample and
records the signal reflected back from the various
acoustic boundaries. Using ultrasonic methods, not only
the travel time of the ultrasonic pulse through the testing
device, and consequently the velocity of waves, but also
the frequency content and the relative energy are
recorded. The ultrasonic inspection test is usually made
in a liquid since the sound waves broadcast in liquid
better than in any other environment. In the manufac-
turing plant, it is observed by ultrasonic inspection that
some pistons have defects. In this study, the causes of
these defects have been investigated. Mainly, two
processes were considered. These are the alphin bath and
the machining process. A variation in the alloying
content is very important in terms of graphitization,
which affects the bonding significantly. It is also thought
that the piston may be damaged during the machining
process. In this research, the effects of variations in the
alloy content and the machining process parameters on
the intermetallic bonding between the diesel piston and
the ring carrier are investigated.
2 EXPERIMENTAL WORK
2.1 Variation in Alloy Content
The alloy contents of non-defective and defective
pistons were investigated statistically. The purpose of the
study is to see the influence of the alloy content on the
intermetallic bonding quality. They were categorized in
two groups, as Group 1 and Group 2, respectively. Group
1 represents the alloy content of the defective pistons and
Group 2 shows the non-defective pistons. A sample
group was selected in both groups and the average and
standard deviation values of the alloys were calculated.
These calculations are summarized in Table 1. The last
column in the table indicates the desired target values for
each individual alloy. Finally, the alloying contents were
related to the intermetallic bonding quality. In addition to
this statistical study, the effect of the graphite flake’s size
on the bonding quality was studied. It was mentioned
previously that some elements accelerate or retard the
graphitization.
Table 1: Statistical study on variations in the alloy content (in mass
fractions, w/%)
Alloying
Elements
Group 1 Group 2
Target
Standard
DeviationAverage
Standard
Deviation
(SD)
Average
Standard
Deviation
(SD)
C 2.61 0.15 2.67 0.08 0.03
Si 2.26 0.12 2.22 0.08 0.03
S 0.03 0.005 0.03 0.01 0.05
P 0.09 0.08 0.08 0.01 0.03
Ni 15.99 0.76 14.84 0.32 0.30
Cr 1.48 0.05 1.24 0.04 0.05
Cu 6.04 0.27 6.34 0.28 0.20
A. F. ACAR ET AL.: EFFECTS OF VARIATIONS IN ALLOY CONTENT AND MACHINING PARAMETERS
392 Materiali in tehnologije / Materials and technology 44 (2010) 6, 391–395
2.2 Effects of the Machining Parameters
For the study of the effects of the machining para-
meters, non-defective pistons were chosen. The aim was
to determine the effects of machining the parameters on
the failure accumulation. The non-defective pistons were
determined by ultrasonic inspection. For this inspection
the piston was immersed into the liquid. If there is any
disconnection in the ring carrier region, the transducer
sound waves make an echo and the waves cannot return
to the receiver. This is an indication of a failure. If the
transducer sound waves are caught by the receiver, it is
an indication of there being no defect. In addition, six of
the samples were also cut and examined. An example of
a non-defective ring carrier is given in Figure 1. Based
on the piston diameters, the machining parameters, such
as the cutting speed and the feed rate of the turning
machines, are determined as given in Table 2. The effect
of the blunt cutter on the defect accumulation was also
investigated.
Table 2: Chip thickness and feed rate for the different piston diame-
ters
Diameters
(mm)
Cutting Speed
r/min
Feeding Rate
mm/min
60−80 1900 0.28-0.40
80−100 1500
100−120 1220
120−140 1050
140−160 900
3 RESULTS AND DISCUSSION
The strength of the intermetallic bonding is related to
the bonding quality. Figure 2 shows the microstructure
of the ring carrier. The left side of the dark boundary
belongs to the Niresist ring carrier material; the right
side shows the Al alloy piston’s microstructure. The
middle dark region displays the debonding area. Various
other structures are presented in Figure 3–5. Figure 3
and 4 show insufficient and inhomogeneous Si thinning,
respectively. It is very important that the Al alloy should
be in an appropriate microstructure for a good-quality
bond. Figure 5 indicates sufficient and homogenous Si
thinning, which is necessary for a good bond. These
microstructures are given for the evaluation and cali-
bration purposes of the bonded region. A defective ring
carrier is shown in Figure 6. The dark region in Figure 6
indicates the defect. It is clear that the ring carrier is
weakly bonded to the Al alloy piston. It this study, the
basic effects of the variation in the alloying content and
the machining parameters are investigated. Table 1
clearly indicates that the standard deviation of the
alloying element with respect to the desired target value
has a great influence on the bond quality. The standard
deviation for Group 2 is close to the target value and a
better bond quality is achieved. The pistons in Group 1
were defective due to the excessive variations in alloy
content. The statistical results explain that by reaching
the target deviation value, the variations in alloy content
affected the bonding in a positive way. It is clear that
choosing the right amount of alloys prevents graphite
formation. However, in addition to the statistical study,
A. F. ACAR ET AL.: EFFECTS OF VARIATIONS IN ALLOY CONTENT AND MACHINING PARAMETERS
Materiali in tehnologije / Materials and technology 44 (2010) 6, 391–395 393
Figure 3: Microstructure of ring carrier (Insufficient Si thinning)
(100X)
Slika 3: Mikrostruktura nosilca obro~ka (nezadostno zmanj{anje
vsebnosti silicija (pove~ava 100-kratna)
Figure 1: Non-defective ring carrier
Slika 1: Nosilec obro~ka brez napake
Figure 2: Microstructure of ring carrier (100X)
Slika 2: Mikrostruktura nosilca obro~ka (pove~ava 100-kratna)
the effect of graphite size on the bonding quality is also
investigated. The experimental observations show that
graphite flakes accumulate on the nonferrous material
and fill the cavities. Eventually, they make the bond
strength weaker due to insufficient bonding surfaces,
since the aluminium alloy does not wet the graphite.
They are seen on the nonferrous material as dark regions
and help in the propagation of the cracks. These findings
indicate that the amount of graphite is very important
and is directly related to the quality of the bond, which
was previously reported 6. It is concluded that the
material that has graphite was dissolved in the alfin bath
and changed the condition of the alfin bath adversely.
These graphite flakes decrease the bonding force. The
strengthening of the intermetallic bonding between the
ferrous and the nonferrous metal through the elimination
of the graphite flakes as a phase is present in the
intermetallic bond region. This can be accomplished by
removing the graphite from the intermetallic bond region
by either removing the graphitic contaminants from the
surface region of the ferrous metal or by sealing the
graphite into the surface of the ferrous metal 6. The
results show that the size of the graphite flakes in the
ring carrier material has an influence on the bonding.
The small size flakes mean a better bond quality. The
recommended flake size and the distribution of the
graphite are shown in Figures 7a–d 7. These figures are
similar to the ASTM flake graphite types 8. These
microstructures can be achieved by the precision of the
alloying elements. The absence of graphite flakes as a
participant in the bond results in an increased bond
strength. In the intermetallic bonding layer, debonding
can be seen at the nonferrous material, which has less
graphite accumulation or attack. The graphite generally
accumulates particle by particle on the nonferrous
material and fills in the moving cavities. This process
reduces the bonding force. The graphite flakes can be
seen as dark regions in the microstructure view and
cause easy crack propagation. The defect during the
formation of the intermetallic bonding depends on the
temperature of the alfin bath and the pressure of the
casting. These small flakes or particles are cooled at a
A. F. ACAR ET AL.: EFFECTS OF VARIATIONS IN ALLOY CONTENT AND MACHINING PARAMETERS
394 Materiali in tehnologije / Materials and technology 44 (2010) 6, 391–395
Figure 7: (a) d = 0.06–0.12, (b) d = 0.03–0.06, (c) d = 0.015–0.03, (d)
d < 0.015 7
Figure 5: Microstructure of ring carrier (Sufficient and homogeneous
Si thinning) (100X)
Slika 5: Mikrostruktura nosilca obro~ka (zadostno in homogeno legi-
ranje silicija) (pove~ava 100-kratna)
Figure 6: Defective ring carrier
Slika 6: Nosilec obro~ka z napako
Figure 4: Microstructure of ring carrier (Inhomogeneous Si thinning)
(100X)
Slika 4: Mikrostruktura nosilca obro~ka (nehomogeno zmanj{anje
vsebnosti silicija (pove~ava 100-kratna)
different stage and show up on the top surface of the iron
part like small- and large-shaped graphite crystal sedi-
ments. As the size of the graphite gets smaller, a stronger
intermetallic bonding is observed. A strong bonding
microstructure is given in Figure 8. The figure reveals
the presence of graphite flakes, extending from the ring
carrier through the bond and into the aluminium alloy.
In the second part of the study, non-defective pistons
were machined by turning at prescribed cutting speeds
and rates. Table 2 shows all the tested conditions. The
machining calculations are summarized in Table 3. All
the possible chip thickness and feed rates were applied.
The cutting forces and powers were calculated for each
case. The selected pistons were inspected by an ultra-
sonic technique for defects. No defects on the piston
were observed from all these studies. Finally, the
machining was performed with a blunt cutter. No defects
were determined to be caused by the blunt cutter. It was
concluded that the machining parameters have no
influence on the intermetallic bonding defects.
4 CONCLUSION
The effects of variations in the alloy contents and the
machining parameters on the strength of the intermetallic
bonding between a diesel piston and a ring carrier have
been investigated. It was determined that the variations
in the alloy contents and the graphite sizes have a great
influence on the bonding quality. As the size of the
graphite gets smaller, a stronger intermetallic bonding is
observed. The graphite accumulation can be resolved by
choosing the right amount of alloying elements. It should
be noted that no defects were caused by the machining-
processes parameters.
Acknowledgement
MOPISAN Piston, Liner, Ring Company in Konya,
Turkey is gratefully acknowledged for their support of
the investigation.
5 REFERENCES
1 Cole, G S., Sherman, A. M., Light weight materials for automotive
applications, Material Characterization, 35 (1995) 1, 3–9
2 Uthayakumar, M., Prabhaharan, G., Aravindan S., Sivaprasad, J.V.,
Study on aluminium alloy piston reinforced with cast iron insert,
International Journal of Material Science, 3 (2008) 1, 1–10
3 Diesel engine tribology, Tribology handbook, CRC Pres LLC, 2001
4 Viala, V.C., Peronnet, M., Bosselet F., Bouix, J., Interface chemistry
in aluminium alloy with iron base inserts, Composites, Part A, 33
(2002), 1417–1420
5 United States Patent, Patent Number: 2,119,833, Date of Patent: June
7, 1938
6 United States Patent, Patent Number: 6,127,046, Date of Patent: Oct
3, 2000
7 Mopisan’s Customer Reference Documents, 2004
8 Radzikowska, J. M., Metallography and Microstructures of Cast
Iron, Metallography and Microstructures, Vol 9, ASM Handbook,
ASM International, 2004, 565–587
A. F. ACAR ET AL.: EFFECTS OF VARIATIONS IN ALLOY CONTENT AND MACHINING PARAMETERS
Materiali in tehnologije / Materials and technology 44 (2010) 6, 391–395 395
Figure 8: Strong bonding microstructure (100X)
Slika 8: Mikrostruktura trdne zveze
Table 3: Effects of machining parameters
Chip
Thickness
mm
Feed
Speed
mm/min
Chip
Volume
mm3
Chip
Shrinkage
Factor
Cutting
Force
(N)
Power
(kW)
2.5 0.28 329.50 75.36 1360 6.46
5.0 0.28 659.40 37.68 3305 32.75
2.5 0.35 412.12 75.36 1360 8.98
5.0 0.35 824.25 37.68 3305 43.67

M. TASI] et al.: ANALYSIS OF THE PROCESS OF CRYSTALLIZATION OF CONTINUOUS CAST ...
ANALYSIS OF THE PROCESS OF CRYSTALLIZATION
OF CONTINUOUS CAST SPECIAL BRASS ALLOYS
WITH THE ACOUSTIC EMISSION
ANALIZA PROCESA KRISTALIZACIJE KONTINUIRNO LITE
SPECIALNE MEDENINE Z METODO AKUSTI^NE EMISIJE
Milo{ Tasi}1, Zoran An|i}1, Aleksandar Vujovi}1, Predrag Jovani}2
1Scientific Research Center, Veliki Park 1, 31000 U`ice, Serbia
2Institute for Multidisciplinary Research, Kneza Vi{eslava 1, 11030 Belgrade, Serbia
pedjaj@ikomline.net
Prejem rokopisa – received: 2010-04-19; sprejem za objavo – accepted for publication: 2010-08-16
Analyses of possibilities of monitoring the crystallization process of continuously cast special brass alloys with acoustic
emission and for establishing a correlation between the microstructure and the recorded acoustic emission signals. With
appropriate selection of parameters for gravitational casting process, continuous casting was performed and samples with a
macrostructure typical of continuous casting were obtained. A laboratory plant for the simulation of the continuous casting and
for the analysis of the crystallization process with acoustic emission was designed. Different energy levels in samples with
different macrostructure, as well as in the defective and non-defective samples, were observed. Two types of sources of signals
were defined: the signal during solidification of correct crystallization and macrostructure of continuous casting and acoustic
emission signal during solidification of samples in with flaws. To check the obtained results, after completion of the
crystallization process, the samples were submitted to external with mechanical loading. The acoustic activity by loading is in
accordance with the results of on-line monitoring of the crystallization process with acoustic emission. The results obtained
show that it is possible to use the acoustic emission for monitoring the crystallization process by continuous and gravitational
casting.
Key words: continuous casting, crystallization, non-destructive examination, acoustic emission
Raziskava je namenjena mo`nosti za spremljanje procesa kristalizacije kontinuirne lite posebne medenine z metodo akusti~ne
emisije in opredelitvi povezave med makrostrukturo in zabele`enimi signali akusti~ne emisije. S primerno izbiro parametrov
procesa gravitacijskega litja je bilo izvr{eno kontinuirno litje in dobljeni so bili vzorci s tipi~no makrostrukturo tega litja.
Pripravljen je bil tudi na~rt za napravo za laboratorijsko simulacijo kontinuirnega litja in analizo procesa kristalizacije z
akusti~no emisijo. Opa`eni so bili razli~ni nivoji energije v vzorcih za razli~no makrostrukturo in vzorci brez napak in z njimi.
Rezultat raziskave je bila definicija dveh izvirov signalov akusti~ne emisije: signal med strjevanjem s pravilno kristalizacijo in
makrostrukturo in signal med strjevanjem vzorcev, v katerih nastajajo napake. Zaradi preverjanja dobljenih rezultatov so bili
vzorci po kon~ani kristalizaciji izpostavljeni zunanji indukciji makrostrukture z mehansko obremenitvijo. Akusti~na emisija
med to obremenitvijo je skladna z rezultati spremljanja procesa kristalizacije z akusti~no emisijo. Sklep je, da je mogo~a
uporaba akusti~ne emisije za spremljanje kristalizacije pri neprekinjenem gravitacijskem litju.
Klju~ne besede: kontinirno litje, kristalizacija, neporu{na preiskava, akusti~na emisija
1 INTRODUCTION
The quantification of microstructures and of their
generating mechanism is a problem that every researcher
in the area of materials needs to deal with. The problem
becomes even more complicated when it is necessary to
define the generated structure of the material without the
application of destructive methods.
The field of non-destructive research provides
different techniques that are used, or may be used, for
controlling the materials quality, such as: vibration
analysis 1, thermography 1,2, X-ray analysis 2, ultrasonic
examination 3,4 and acoustic emission 5,6.
The principle of the acoustic emission is that, due to
changes occurring in the material, there is a sudden
release of accumulated deformation energy, in form of
mechanical waves that are transmitted through the
material, and detected by sensors. The phenomenon of
generating mechanical waves with release of a part of the
deformation energy in the material is called acoustic
emission 7,8.
Given that all the materials used for construction
purposes (metals, alloys, glass, ceramics, wood, concrete
and polymers) produce, under certain conditions,
acoustic emission signals, this method can be very
successfully used: 8,9,10
– in tensile hardness tests,
– for structural composition and material characte-
ristics control,
– in phase transformations in material tests,
– for controlling vessels and water-pipelines,
– for controlling aircraft and spacecraft constructions,
– in welded tacks tests,
– in crack tests,
– in material fatigue monitoring,
– in monitoring a crack’s development at low tempe-
ratures,
– in material solidification monitoring, which is still in
its infancy 11,12.
Materiali in tehnologije / Materials and technology 44 (2010) 6, 397–401 397
UDK 669.35.018:621.745:546.821 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(6)397(2010)
In this study, the main objective was to define a
method for monitoring the process of continuous casting
that might effect the casting quality and enable the
monitoring and management. Since this is a continuous
process, conventional monitoring methods with
destruction are inadequate because the process needs to
be decelerated during each sampling. In addition to this,
monitoring performed in this way requires much time,
labour and material, with the ever-present problem of the
casting quality of two samplings. The results of the
analysis of the process of crystallization of the
continuous casting of brass-type special alloys with
acoustic emission are shown here. Every change in the
level of the acoustic emission signals characterises a
corresponding structure object, which is later identified
in a metallographic analysis. Thus, a particular moni-
toring diagram "structure – the acoustic emission signal"
is defined, which may be used for the on-line characte-
risation of a material during the process of continuous
casting.
2 EXPERIMENTAL
The material used in this research belongs to the
group of the special brass alloy type CuZn21Al2As.
Two groups of experiments were performed. The
objective of the first group was to record the generating
of the "ideal" microstructure in terms of acoustic activity,
i.e., to determine the belt of acoustic emission signals
that guarantees the correct crystallization and the
obtaining the continuous casting macrostructure. The
second group of experiments included the recording of
acoustic emission signals of samples with artificially
induced flaws with inadequate cooling regime and the
inflow of molten metal.
The experiments comprised the physical stimulation
of continuous casting by gravitational casting, the
defining the characteristic signals that appear in the
processes of solidification and further cooling. A
continuous casting simulation combining gravitational
casting and changes in velocity of the casting mould may
be observed in time and ways of cooling of the casting
mould, so that the macro-effect on the mould is the same
as if it has gone through the cooling zones of continuous
casting. For this reason, an appropriate steel casting
mould was designed and produced, which, under specific
cooling conditions, has similar velocities of conductivity
and heat outlet to continuous casting. Figure 1 shows the
scheme of the experimental casting mould.
The casting mould has sensor holders and thermopair
for measuring the temperature during cooling attached.
The acoustic emission sensor holder, shown in Figure 2,
is specially designed for the work at increased tem-
perature. It is made in the form of a ribbed tube of an
aluminium alloy and with the aim to prevent thermal
damage to the sensor. In its middle part, due to the
identity of the material whose acoustic activity is being
examined and the material which the signal passes
through to the sensor, there is a brass bar that functions
as a signal conductor. The sensor is a piezoelectric tile,
type BMF 10P-5, with a resonance frequency of 180kHz.
The total amplification of the acoustic emission signals
is different, from 5000 to 10000. The frequency response
of the amplifying level from 100kHz to 3MHz is limited
by the use of a filter, which does not affect the total
amplification and the resonance field of the sensor. The
processing of the acoustic emission signals is performed
by the method of oscillation counting recorded by the
sensor (the ring-down counting method), which is based
on counting each outthrow at the decision level by the
signal obtained at the sensor exit. Every time that level is
outthrown, it generates a signal registered by the gauge.
The total number of signals surpassing the specific (set
up) decision level, appears at the unit exi as an analogous
signal. This signal is written on the plotter, via an A/D
convertor, or a PC. During the research the RMS method
(effective value of the voltage) for analysing the acoustic
emission signals during the measurement interval, by
M. TASI] et al.: ANALYSIS OF THE PROCESS OF CRYSTALLIZATION OF CONTINUOUS CAST ...
398 Materiali in tehnologije / Materials and technology 44 (2010) 6, 397–401
Figure 2: The scheme of the acoustic emission sensor holder
Slika 2: Shema nosilca senzorja akusti~ne emisije
Figure 1: The scheme of the experimental casting mould
Slika 1: Shema eksperimentalne livne kokile
which, applying the appropriate transformations, the
original signal is transformed into a signal whose shape
is much clearer for reading.
Figure 3 shows the laboratory plant for monitoring
the acoustic emission signals. After recording the signals
of the acoustic emission, the obtained samples are
submitted to the quality control of the macrostructural
properties. Finally, in order to check the results of the
research, after the completion of the solidification
process, the samples are submitted to external induction
by mechanical loading.
3 RESULTS AND DISCUSSION
The quality control of the macrostructural properties
indicates that with an appropriate choice of parameters
of gravitational casting into a designed metal casting
mould (a method of pouring molten metal, cooling
velocity, etc.), it is possible to simulate the continuous
casting and obtain castings with a continuous casting
macrostructure.
Figure 4 shows the macrostructure of samples with
correct crystallization and continuous casting macro-
structure, whereas Figure 5 shows the recording of the
acoustic emission signal for the observed sample.
Analyzing the recording of acoustic emission signal
with correct crystallization and continuous casting
macrostructure, it can be concluded that there is a sudden
increase in the acoustic activity immediately after the
pouring of molten metal into the casting mould, which is
the result of the primary solidification of the -solid
solution and the occurrence of friction between the solid
metal and the casting mould. The next phase is the phase
of the linear growth of acoustic activity until the sample
is completely solidified. In addition, it is evident that
there is a uniform distribution of the effective (RMS)
voltage, which indicates the absolutely correct crystalli-
zation and the obtaining of samples without flaws. This
is illustrated by the recording of the macrostructure of
the analysed sample.
Figure 6 shows the macrostructure and Figure 7
shows the acoustic emission signal of samples with a
flaw.
There is a more intensive acoustic activity in samples
with any type of flaw and gravitational casting
macrostructure (Figure 6) during the crystallization
process (Figure 7) in comparison to the acoustic activity
of samples with the macrostructure of continuous
casting. This results from the fact that the friction
between the small grains characteristic of gravitational
casting, due to the larger total area, is higher than the
friction among large grains characteristic of continuous
casting. A more intensive acoustic activity in samples
with macrostructure of gravitational casting can be
associated with the fact that the grain boundaries
represent an area of a disturbed crystal structure and that
these disorders are more distinctive if there is a greater
difference in orientation of neighbouring crystals. In
addition, a small-grained microstructure may be caused
by effects preventing grain enlargement, and all this
initiates more acoustic activity in comparison to the
M. TASI] et al.: ANALYSIS OF THE PROCESS OF CRYSTALLIZATION OF CONTINUOUS CAST ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 397–401 399
Figure 4: The scheme of the macrostructure of the sample with
absolutely correct crystallization and a continuous casting macro-
structure
Slika 4: Shema makrostrukture kontinuirno litega vzorca s pravilno
kristalizacijo
Figure 3: The laboratory plant for monitoring acoustic emission
signals during solidification
Slika 3: Laboratorijske naprava za bele`enje signalov akusti~ne
emisije med strjevanjem
Figure 5: The recording of the acoustic emission signal for the sample
with an absolutely correct crystallization and continuous casting
macrostructure
Slika 5: Registracija signala akusti~ne emisije za vzorec s pravilno
kristalizacijo in kontinurne lite makrostrukture
acoustic activity in samples with continuous casting
macrostructure.
When analysing the results of the completed
examination, it is evident that all the samples show
acoustic activity during the crystallization process. In
samples with out flaws in the latter phase (Figure 4), the
analysis of the RMS voltage of acoustic emission signals
(Figure 5) does not indicate the existence of any
characteristic signal after the crystallization signal.
While monitoring the acoustic activity, there is a specific
even distribution of RMS voltage, without any distinctive
peaks. There is also a specific belt of values of the
acoustic emission signals, which, for a particular mould,
guarantees a correct crystallization and the obtaining of
the continuous casting macrostructure, where a three-
zone distribution of crystals (from the small globular and
column-like crystals to the large globular ones in the
central part of the crystals) are clearly visible. If the
signal comes out of this belt, the crystallization is
irregular. While monitoring the acoustic activity of
samples with flaws of any type (Figure 6), the
distribution of the RMS voltage is irregular, with clearly
visible peaks (Figure 7).
The obtained results show that it is possible to use the
acoustic emission method for monitoring the crystalliza-
tion process during continuous casting. Two types of
signal sources are defined here:
– the acoustic emission signal during solidification of
samples with correct crystallization and continuous
casting macrostructure,
– the acoustic emission signal during solidification of
samples with flaw (crack, inclusion, etc.) and gra-
vitational casting macrostructure.
With the aim to check the results of the research, the
formed (cooled) samples are submitted to changes of the
macrostructure with mechanical loading. With changes
of structure a specific material activity, i.e., the existence
of "frozen" faults in the material is detected.
Figure 8 shows the acoustic activity of loaded
sample with an absolutely correct crystallization and
continuous casting macrostructure.
M. TASI] et al.: ANALYSIS OF THE PROCESS OF CRYSTALLIZATION OF CONTINUOUS CAST ...
400 Materiali in tehnologije / Materials and technology 44 (2010) 6, 397–401
Figure 8: The acoustic activity of the externally induced sample with
completely correct crystallization and a continuous casting macro-
structure
Slika 8: Akusti~na emisija obremenjenega vzorca s pravilno kristali-
zacijo in makrostrukturo kontinuirnega litja
Figure 7: The acoustic emission signal for the sample with a flaw
Slika 7: Akusti~na emisija vzorca z napako
Figure 6: The macrostructure of the sample with a flaw
Slika 6: Makrostruktura vzorca z napako
Figure 9: The acoustic activity of an externally induced sample with a
flaw
Slika 9: Akusti~na aktivnost zunanje obremenjenega vzorca z napako
On samples without flaws in macrostructure, a
uniform activity increase under the mechanical load is
evident, which means that the voltages in the material are
uniformly distributed in the volume. There was no
particular inhomogeneity of the material, and the
analysis of the crystallization process, which leads to
that particular structural stabilization, is adequately
recorded by the acoustic emission signal.
Figure 9 shows the acoustic activity of loaded
sample with a flaw.
The recording of the acoustic activity indicates that
during loading of the material with a flaw there are
dramatic amplitude oscillations, which means that the
voltages in the material are not uniformly distributed in
the volume. According to the results of the monitoring of
the process of crystallization this is characteristic for
defective samples according to the acoustic emission
method.
4 CONCLUSION
With an appropriate selection of parameters of
gravitational casting into a metal casting mould (the style
of pouring, a definite cooling velocity, etc.), it is possible
to simulate continuous casting and obtain castings with
the continuous casting macrostructure.
The results obtained from the analysis of the process
of solidification indicate that it is possible to use acoustic
emission as a method for recording and for monitoring
the crystallization process during continuous and gravity
casting.
Two types of sources of the acoustic emission signals
were defined:
the acoustic emission signal during the solidification
of samples with a completely correct crystallization and
a continuous casting macrostructure,
the acoustic emission signal during the solidification
of samples with of flaws (crack, inclusion, etc.) and
gravitational casting macrostructure.
All the signals of the acoustic emission show acoustic
activity at the beginning of the monitoring process, i.e.,
the acoustic activity of the crystallization process. In
samples with a correct crystallization, the RMS analysis
does not indicate the existence of any characteristic
signal after the crystallization signal. There is also a
specific belt of acoustic emission signals, which
guarantees completely correct crystallization for that
particular sample. If the signal originates out of this belt,
the sample has a flaw.
By inducing the structure with a mechanical load in
both samples, defective and non-defective, the acoustic
activity is identified, which is in accordance with the
results of the on-line monitoring of the crystallization
process using the acoustic emission method.
Further research will be aimed to define the
dependence of the level of acoustic emission signal on
the type and the location of the flaw in the sample on the
basis of the analysis of the values of the RMS voltage,
the development of the adaptive system for the regulation
of the process of continuous casting on the basis of the
developed on-line method, as well as on the application
of the obtained results of other casting technologies.
5 REFERENCES
1 Hameeda Z., Honga Y. S., Choa Y. M., Ahnb S. H., Songc C. K.,
Condition monitoring and fault detection of wind turbines and
related algorithms: A review Renewable and Sustainable Energy
Reviews, 2007
2 Handbook of composites. Edited by S. T. Peters. Publisher in 1998
by Chapman & Hall, London, 839–855
3 Rai{utis R., Ka`ys R., Ma`eika L., Application of the ultrasonic
pulse-echo technique for quality control of the multi-layered plastic
materials, NDT&E International, 41 (2008) 300-311
4 Raghavan A., Carlos E., Cesnik S., Review of guided-wave structural
health onitoring, The Shock and Vibration Digest., 39 (2007),
91–114
5 Anastassopoulos A. A., Kouroussis D. A., Nikolaidis V. N., Struc-
tural integrity evaluation of wind turbine blades using pattern
recognition analysis on acoustic emission data, 25th European
Conference on Acoustic Emission Testing – EWGAE 2002, Prague,
Czech Republic, 11–13 September 2002, 1–8
6 R. Rai{utis, E. Jasiûnienë, R. [literis, A. Vladi{auskas, The review of
non-destructive testing techniques suitable for inspection of the wind
turbine blades, ULTRASOUND, 63 (2008) 1, 26–30
7 Purvis L. A., Kannatey-Ashibu E., Pehlke R. D.: Evaluation of
Acoustic Enission from Sand Cast Aluminium Alloy 319 During
Solidification and Formation of Casting Defects – Department of
Materials Science and Engineering Department of Mechanical The
University of Michigan Ann Arbor, Michigan, 1990, 31
8 D. Mitrakovi}, D. Sc. Thesis, Faculty of Technology and Metallurgy,
University of Belgrade, 1984 (in Serbian)
9 N. A. Bunina: Issledovanie plasti~eskoi deformacii matallov meto-
dom akusti~eskoi emissii – Leningradskii in`enerno-ekonomi~eskii
institut, Leningrad, 1990, 30–58
10 C. R. Heiple, S. H. Carpenter: Acoustic emission produced by the
delta-to-alpha phase transformation in Pu-Ga alloys, Metallurgical
Transactions A, 23A (1992) 3, 779–783
11 M. Tasi}, D. Sc. Thesis, Faculty of Technology and Metallurgy,
University of Belgrade, 1997. (in Serbian)
12 M. Tasi}, Z. An|i}, D. Bojovi}, Contribution for cristallization
process analysis of continuous casting by acoustic emission method,
Second International Conference on Materials and Manufacturing
Technologies, Matehn '98, Proceedings, Vol. 2, Cluj-Napoca,
Romania, 1998, 879–884
ACKNOWLEDGEMENT
The authors wish to thank Air. Institute of Ribinsk,
Russia, for help with the instrumental techniques and the
interpretation of the results.
M. TASI] et al.: ANALYSIS OF THE PROCESS OF CRYSTALLIZATION OF CONTINUOUS CAST ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 397–401 401

D. SMOLOVI] et al.: THE INFLUENCE OF ADDING EMULSION FLOCCULANTS ...
THE INFLUENCE OF ADDING EMULSION
FLOCCULANTS ON THE EFFECTS OF RED-MUD
SEDIMENTATION
VPLIV DODATKA EMULZIJSKIH FLOKULANTOV NA POJAVE
PRI SEDIMENTACIJI RDE^EGA BLATA
Dejan Smolovi}1, Mira Vuk~evi}2, Dragoljub Ble~i}2
1KAP – Alumina plant, 81000 Podgorica, Montenegro
2Faculty of Metallurgy and Technology, Cetinjski put b.b., 81000 Podgorica, Montenegro
mirav@ac.me
Prejem rokopisa – received: 2010-06-22; sprejem za objavo – accepted for publication: 2010-09-21
The main objective of this investigation was to define, based on an industrial probe, the influence of emulsion flocculants on the
effects of red-mud precipitation. The precipitation velocity of the red mud as well as the transparency of the liquid phase create
the basis for the comparison of the efficiency between the conventional flocculants and the emulsion flocculants. The
investigation was also focused on the characteristics of the precipitated mud presented as the content of dry particles as well as
the granulation of red mud. The influence of the suspension viscosity on the loading of the mixer inside the decantation vessels
was also investigated. All of the precipitation experiments were carried out within the automatic Cytec unit for the preparation
and dosage of the liquid flocculants. The parallel investigation of the efficiency of two different flocculants was performed in
separate decantation vessels.
The results show that the red-mud precipitation velocity during the experiment with the emulsion flocculants is several time
higher than the precipitation velocity obtained with synthetic flocculants. The emulsion flocculants also enable a better solution
transparency with a clear and distinct boundary between the two phases (liquid and solid). Besides that, the consequence of the
emulsion flocculants’ dosage is a higher content of dry particles in the suspension as well as an increased granulation and a
lower viscosity. Using emulsion flocculants increased the decantation vessels’ productivity, improved the quality of the
aluminate’s solution and improved the effects of red-mud rinsing.
Key words: flocculants, red mud, precipitation velocity, solution transparency, aluminate’s solution, decantation vessel
Glavni cilj te raziskave je bil dolo~itev vpliva emulzijskih flokulantov na pojave obarjanja rde~ega blata na podlagi industrijskih
preizkusov. Hitrost obarjanja rde~ega blata in transparenca teko~e faze sta osnovi za primerjavo u~inkovitosti med
konvencionalnimi in emulzijskimi flokulanti. Cilj raziskave so bile tudi zna~ilnosti izlo~enega blata pa tudi vsebnost suhih
delcev in granulacija rde~ega blata. Vpliv viskoznosti suspenzije na obremenitev me{alnika v posodi za dekantacijo je bil tudi
dolo~en. Vse preizkuse smo izvr{ili v avtomatski napravi Cytec za pripravo in doziranje teko~ih flokulantov. Vzporedne
preiskave u~inkovitosti razli~nih flokulantov so bile izvr{ene v lo~enih posodah za dekantacijo. Rezultati ka`ejo, da je hitrost
obarjanja rde~ega blata med preizkusom z emulzijskim flokulatom ve~krat ve~ja od hitrosti pri uporabi sinteti~nih flokulantov.
Ti flokulanti zagotavljajo tudi bolj{o transparenco emulzije z jasno in razlo~no mejo med obema fazama (teko~a in trdna).
Posledica doziranja emulzijskih floukantov so tudi ve~ja vsebnost suhih delcev v suspenziji, pove~ana granulacija in ni`ja
viskoznost. Uporaba emulzijskih flokulantov pove~a produktivnost posod za dekantacijo, pove~a kakovost aluminatne raztopine
in izbolj{a izpiranje vplivov rde~ega blata.
Klju~ne besede: flokulanti, rde~e blato, hitrost obarjanja, transparenca raztopine, aluminatna raztopina, posoda za dekantacijo
1 INTRODUCTION
One of the most important operations in the Bayer
process of alumina production is the separation of red
mud from the aluminate’s solution in decantation vessels
using flocculants of different types. The separation
operation is not based on the chemical reactions, but on
the hydrodynamic conditions and the system’s characte-
ristics, the separation conditions, the type of flocculants
used and the construction of the decantation vessel1,2.
The mineralogical content of bauxite, the dissolution
conditions as well as the red-mud granulation all have an
influence on the red-mud precipitation 3.
The main demand from the industrial Bayer process
is the final result in the form of relatively clean alumina.
The main precursor for the process is the aluminate’s
solution without any traces of impurities. The precipi-
tation of solid particles is often defined by Stocks’
equation, approved by Richardson and Zaki (1954).
According to this equation, the velocity of a particle’s
precipitation is directly proportional to the square
exponent of the aggregate’s size and so a small increase
in the particle size has a significant influence on the
precipitation velocity. This is particularly visible for
small particles ( 50 μm) with the very long preci-
pitation time. The process of particle aggregation is of
great importance for the separation of red mud from the
aluminate’s solution. The dosage of the flocculants
facilitates the creation of red-mud aggregates and the
improved possibility of gravitation decantation 5.
The preliminary use of common organic compounds
as the precipitation facilitators was replaced with
synthetic poly-acrylates because of the improved values
of the decantation velocities. On the other hand, the
Materiali in tehnologije / Materials and technology 44 (2010) 6, 403–406 403
UDK 669.715:628.3 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 44(6)403(2010)
transparency of the final liquid phase was lower in
comparison to the effects of the organic flocculants.
Nowadays, a new class of synthetic hydroxamathic
poly-acryl-amides flocculants has been developed. The
aim was to obtain improved red-mud precipitation
velocities and improved transparency of the liquid phase.
The objective of this investigation was to compare
the two types of flocculants’ efficiency by using the
hydroxamatic poly-acryl-amides and the conventional
synthetic flocculants used in the alumina factory,
Aluminum Plant, Podgorica, Montenegro.
2 METHODS OF WORK
An investigation of the emulsion flocculants’ effi-
ciency as well as the influence on the effects of red-mud
precipitation started with the industrial low-con-
centration suspension that has the following content
(Na2Ok) = 160.72 g/L; k = 1.56, as well as the
following content of red mud:
w (Al2O3) = 15.3 %
w (F2O3) = 41.8 %
w (Na2Or) = 0.341 %
w (Na2Ouk) = 5.42 %
w (SiO2) = 13.17 %
w (Zn) = 0.0307 %
w (TiO2) = 5.13 %
w (CaO) = 2.67 %
k = 4.495
The experimental conditions were as follows: T =
95–100 °C;  = 1320–1340 g/L.
The preparation and dosage of the emulsion
flocculants in industrial conditions was carried out in the
Cytec automatic unit. The emulsion flocculants were
introduced into the separate decantation vessels in two
focal points. The gravitation dosage into the decantation
vessels was used through the separation box filled with
the low-concentration solution after the leaching of the
bauxite. The first focal point was on the dosage ring of
the decantation vessel and the second one was on the
dosage tube through which the suspension is pumped
into the dosage ring of the decantation vessel. The
preparation and the dosage of the synthetic flocculants in
the second decantation vessel were carried out in the
existing equipment for the preparation and dosage of the
same flocculant. The synthetic flocculants were also
introduced through the two points at the same positions
as the emulsion one. In some of the experimental parts
we even introduced the third flocculant, i.e., starch.
The same conditions for the investigation of the
precipitation velocity were used for all three different
flocculants (the velocity was measured in the 1L vessel
in all of the cases). The used flocculants were as follows:
the natural one (starch) (Ipokol Eg 720), the synthetic
flocculants (anion type A-185 HMW) as well as the
hydroxamatic poly-acryl-amides (HXPAM) HX 300 as
the representative of the emulsion flocculants. All of the
flocculants were prepared in the same conditions as the
water solution with a low content of  (Na2Ok)  50
mg/L and under a temperature of 50 °C.
The concentration of starch (Ipokol EG 720) was
1 %, with a dosage level of 1.5 kg/t red mud; the con-
centration of synthetic flocculants (A-185 HMW) was
0.05 %, with a dosage level of 0.08 kg/t red mud; and the
concentration of the emulsion flocculants HX 300 was
0.25 % with a dosage level of 0.45 kg/t red mud.
The transparency of the liquid phase was determined
by its content of dry particles. The content of dry
particles in the red mud was also determined through
probes from the drain valve.
The loading of the mixer within the decantation
vessel in the case of using the emulsion flocculants and
the synthetic flocculants was followed up continuously
with dynamometers on both of the decantation vessels.
3 RESULTS AND DISCUSSION
The red-mud precipitation velocity as a function of
the time during the dosage of different flocculant types is
shown in Figure 1. The results show that the highest
value of the precipitation velocity is obtained after a
period of 1 min, using the HX300 flocculants and it is
D. SMOLOVI] et al.: THE INFLUENCE OF ADDING EMULSION FLOCCULANTS ...
404 Materiali in tehnologije / Materials and technology 44 (2010) 6, 403–406
Figure 2: Solution transparency as a function of flocculants type used
in the process
Slika 2: Transparenca raztopine v odvisnosti od vrste flokulanta upo-
rabljenega v procesu
Figure 1: Variation of the red-mud precipitation velocity as the func-
tion of the flocculants type
Slika 1: Sprememba hitrosti obarjanja rde~ega blata v odvisnosti od
vrste flokulanta
0.96 mm/s. Using A-185 HMW flocculants the highest
level of velocity is reached after 300 s in the range of 0.4
mm/s. In terms of velocity, starch shows the lower level
of velocity (0.2 mm/s) after a period of 300 s.
The dosage of HX 300 flocculants enables the
formation of clearly divided zones: the liquid zone
(aluminate’s solution) and the solid zone (red mud). A
kind of transitional zone is present in the application of
the A-185 HMW flocculants.
The transparency of the liquid phase in all three cases
was measured with an optical prism "Ciba flocculants".
The highest level of transparency on the prism’s scale is
54, which is equivalent to the transparency of water. In
our cases, the best transparency (42) was obtained in the
case of the starch application, and it was lower for the
HX 300 flocculants application (35). The lowest value of
the transparency (20) was obtained with the application
of A-185 HMW flocculants (Figure 2).
The precipitation velocity as a function of time for
different quantities of emulsion flocculants HX 300 is
shown in Figure 3. The results show an increasing trend
for the precipitation velocity with an increased quantity
of the flocculants HX 300. The lowest velocity level was
0.8 mm/s, obtained with a dosage of 500 g/t red mud.
This increases to 1.2 mm/s with a dosage of 600 g/t red
mud. The highest level of velocity (1.5 mm/s) was
obtained with a dosage of 650 g/t red mud.
The liquid-phase transparency was followed using
the content of dry particles in the remaining residue for a
period of 4 months. This is the period when the
industrial probe of the emulsion flocculants is compared
with the effects on the content of dry particles in the
aluminate’s solution in the case of the application of the
flocculants’ A-185 HMW. The content of dry particles in
the liquid phase using A-185 HMw as well as HX 300 is
shown in Figure 4. The results show that the content of
dry particles in the liquid phase, after a dosage of HX
300, is reduced to an average value of 3g/L. The dosage
of A-185 HMW reduced that content to an average value
lower than 1 g/L.
D. SMOLOVI] et al.: THE INFLUENCE OF ADDING EMULSION FLOCCULANTS ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 403–406 405
Figure 4: Content of solid particles in overflow settlers before and
after the application of HX 300 flocculants
Slika 4: Vsebnost trdnih delcev in pretok uporabo in po njej flokulan-
tov HX 300
Figure 5: Dry particles’ content in the red mud before and after the
application of HX 300 flocculants
Slika 5: Vsebnost suhih delcev v rde~em blatu pred uporabo in po njej
flokulantov HX 300
Figure 3: Variation of the red-mud velocity as a function of time
during the application of three different concentrations of flocculants
HX 300
Slika 3: Sprememba hitrosti rde~ega blata v odvisnosti od ~asa pri
uporabi treh razli~nih koncentracij flokulanta HX 300
Figure 7: SEM images of the red-mud samples obtained with the
application of A-185HMW flocculants (a) and HX 300 flocculants (b)
Slika 7: SEM-posnetek vzorca rde~ega blata, nastalega pri uporabi
flokulantov A-185 HMW (a) in HX 300 (b)
Figure 6: Portion of the individual fractions (percentage) in red mud
settled as a function of different flocculants type application
Slika 6: Dele` posameznih sestavin (odstotki) v usedlini rde~ega blata
v odvisnosti od uporabe razli~nih vrst flokulatov
During the industrial probe with the HX 300
flocculants, the content of solid particles in the red mud
was also investigated. The results obtained before and
after the dosage of the HX 300 are shown in Figure 5.
The results show that the content of solid particles in the
red mud increases from an average value of 400 g/L
(A-185 HMW) to 500 g/L after the HX 300 application.
The conclusion is that the red-mud rinsing, as well as the
alkali content in it, is more effective after the HX 300
application. A granulometric analysis of the red-mud
particles obtained during the A-185 HMW and HX 300
application is shown in Figure 6.
The results show that using the HX 300 increases the
content of huge fractions in red mud. The SEM analysis
indicates the formation of bigger, spherical aggregates
(the “closed type”) (Figure 7).
The red-mud suspension’s viscosity was followed
indirectly through the mixer branches’ load torque in the
decantation vessel, with a dosage of A-185 HMW and
HX 300 flocculants respectively. The results of thr
comparison are shown in Figure 8. The dosage of HX
300 decreases the mixer branches’ load torque. This can
be explained by the lower viscosity of the red mud
obtained by using the HX 300 flocculants.
4 CONCLUSIONS
The dosage of HX 300 enables higher sedimentation
velocities of red mud in comparison with the dosage of
the A-185 flocculants. The consequence is an increased
productivity of the decantation vessel and the decreased
number of vessels in action.
The dosage of HX 300 flocculants (in comparison
with the A-185 dosage) enables a distinct boundary
between the solid and the liquid phases. The liquid phase
has a lower content of solid particles; a fact that
improves the filtration conditions on Kelly filters as well
as the transparency of the aluminate’s solution.
Precipitated red mud obtained with the HX 300
dosage (in comparison with A-185 HMW) has an
increased content of solid particles, which improves its
rinsing and alkali loss. The red mud obtained under these
conditions has a bigger granulation caused by the
chemical composition of the flocculants.
In the case of the HX 300 dosage, the red-mud
suspension has a lower viscosity, which improves the
condition of extracting it from the decantation vessel,
prevents the formation of mud deposits on the walls and
the arms of the mixer, and thus lowers the mixer
branches’ load torque and increases the decantation
vessel’s life time.
5 LITERATURE
1 Ryles, R. G., Avotins, P. V., Cytec Industries Inc, Superfloc HX, a
new technology for the alumina industry, Fourth International
Alumina Quality Workshop, Darwin 2–7 June 1996
2 Fitch, B., Thickening theoris an analysis, AlChE Journal, 36 (1993)
1, 27–36
3 Settling characteristics of QAL Red mud., E1445 Thesis, Department
of Chemical Engineering by Toby Marsh, Supervisor: dr Tony
Howes., October, 1998
4 Blecic, D., Adzic, M; Analysis of the results obtained by experi-
mental investigation of aluminate solution from red mud, 6th
Yugoslav International Symposium on Aluminium. I. Bauxites and
Extractive Metallurgy, Yugoslavia, 1990, 113–123
5 Coulson, J. M., Richardson, J. F., Brackhurst, J. F., Harker, J. H.,
Chemical engineering Volume 2, Fourth Edition, Particle technology
and separation processes, Pergamon Press, Oxford, 1991
6 Richardson, J. F., Zaki, W. N, Sedimentation and Fluidisation: Part I,
Trans. Ins. Chem. Eng., 32 (1954) 35
D. SMOLOVI] et al.: THE INFLUENCE OF ADDING EMULSION FLOCCULANTS ...
406 Materiali in tehnologije / Materials and technology 44 (2010) 6, 403–406
Figure 8: Comparison of the mixer branches’ load torque in settlers
during the application of synthetic flocculants (A-185 HMW) and
emulsion flocculants (HX 300)
Slika 8: Primerjava obremenitev propelerjev me{alnika v usedalnikih
pri uporabi sinteti~nih flokulantov (A-185H MW) in emulzijskih
flokulatov (HX 300)
N. [TREKELJ ET AL.: METALLOGRAPHIC CHARACTERIZATION OF THE JOINING PART ...
METALLOGRAPHIC CHARACTERIZATION OF THE
JOINING PART FOR EN 10216-1 AND EN 10083-3 MADE
WITH DIFFERENT CONSUMABLES
METALOGRAFSKA KARAKTERIZACIJA VEZNEGA ^LENA IZ
EN 10216-1 IN EN 10083-3, IZDELANEGA Z RAZLI^NIMI
DODAJNIMI MATERIALI
Neva [trekelj1, Marko Lampi~1, Ladislav Kosec1, Bo{tjan Markoli1
Faculty of Natural Sciences and Engineering, University of Ljubljana, A{ker~eva 12, Slovenia
strekeljn@gmail.com
Prejem rokopisa – received: 2010-03-22; sprejem za objavo – accepted for publication: 2010-08-23
The microstructures of the base materials (EN 10216-1 and EN 10083-3) and three consumables (EVB 50, EVB CrMo and
INOX R 29/9) were investigated in the as-welded and heat-treated states. The base material where the fracture occurred was EN
10216-1, which is why the characterization was focused only on that side of the weld. The microstructure of the base material
EN 10216-1 consisted mainly of ferrite. The weld with EVB 50 showed a ferritic and bainitic microstructure. The weld with the
consumable EVB CrMo showed a microstructure that consisted of ferrite and some bainite. The microstructure of the weld with
the consumable INOX R 29/9 showed an austenic and ferritic microstructure and the fusion line is more pronounced than in the
previous two cases. Heat treatment (1 h, 600 °C) affected only the columnar dendrite-like morphology in the case with the
consumable EVB CrMo. The weldment is free of typical welding defects, such as inclusions, porosity, microcracks, etc. Out of
the three consumables, the EVB CrMo was slightly better in terms of its ultimate tensile strength.
Keywords: microstructure, arc-welding, steel, consumables
Preiskovana je bila izhodna in toplotno obdelana mikrostruktura dveh osnovnih materialov (EN 10216-1 in EN 10083-3) in treh
dodajnih (EVB 50, EVB CrMo in INOX R 29/9). Prelomi so nastopili na strani z osnovnim materialom EN 10216-1, zato je bila
karakterizacija osredinjena samo na to stran zvara. Mikrostruktura osnovnega materiala EN 10216-1 je bila sestavljena
ve~inoma iz ferita. Mikrostruktura vara z dodajnim materialom EVB 50 je bila feritno bainitna, mikrostruktura podro~ja vara z
EVB CrMo je bila prav tako iz ferita in bainita, var z dodajnim materialom INOX R 29/9 pa je izkazoval avstenitno feritno
mikrostrukturo. ^rta spajanja je bila o~itno tanj{a kot v drugih dveh primerih dodajnih materialov. Toplotna obdelava (1 h, 600
°C) je vplivala le na stebrasto dendritno morfologijo v primeru z dodajnim materialom EVB CrMo. Pri vseh treh zvarih ni bilo
najti tipi~nih defektov, kot so vklju~ki, poroznost, mikrorazpoke itd. Med vsemi tremi dodajnimi materiali je v smislu natezne
trdnosti dodajni material EVB CrMo nekoliko bolj{i.
Klju~ne besede: mikrostruktura, ro~nooblo~no varjenje, jeklo, dodajni materiali
1 INTRODUCTION
Welding is a very useful procedure for joining
different materials of different shapes that cannot be
manufactured by casting. It is also used for different
repairs of worn and poorly manufactured pieces. How-
ever, it is very important as to how the welding technique
and the consumables are chosen. The consumable should
be similar to the base material. When welding, the focus
must be on the heat input, which significantly affects the
development and the properties of the weld, i.e., the
microstructure, the appearance of cracks and inclusions.
In some cases a protective gas has to be employed as
well as a pre-welding heat treatment.1,6,8
In this paper the constitution and the development of
the microstructure of the base material and the weld
material were investigated (in the as-welded and
heat-treated states) in the links for drilling rigs using
metallographic methods, i.e., optical microscopy and
scanning electron microscopy with energy-dispersive
X-ray spectroscopy. Hardness measurements were made
with a Vickers testing machine and tensile tests were
conducted for all three consumables.
2 THEORETICAL BACKGROUND
For quality welds in arc welding an awareness of the
metallographic properties of the base and filler materials
is very important. The basic properties of welds are very
dependent on the chemical composition of the base and
filler materials and on the regime of the welding. It is
very important to identify the microstructural ingredients
of the weld and their morphology in connection with the
properties of non-equilibrium solidification, which is
usually present during the welding procedures.1
When welding heavy alloyed steel, such as chromium
and nickel alloyed steel, it is essential for the metallo-
graphic analysis to use a Schäffler diagram.2 With this
one can make an estimate of the presence and the
fraction of austenite, ferrite, bainite and martensite,
depending on the chemical composition.
The most appropriate procedure for welding the links
for the drilling rigs was patented in 1998,3 while the
filler materials were patented in 1996.4 In the literature
the data on the welding parameters is scarce for these
types of welds, which also goes for the microstructural
behavior while welding.1
Materiali in tehnologije / Materials and technology 44 (2010) 6, 407–412 407
UDK 669.14:621.791.05 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 44(6)407(2010)
The section through the weldment is presented in
Figure 1.
3 EXPERIMENTAL
3.1 Materials
The weld was made from two parts: a support made
of EN 10216-1 steel and the flange attached to it made of
EN 10083-3 steel. The chemical compositions of these
two alloys are given in Table 1. For the welding we used
three different consumables: EVB 50, EVB CrMo and
INOX R 29/9 (Table 2). The welding root had a typical
šV’ shape.
The workpiece was made of two parts. The welding
was in the form of arc-welding with three different
consumables, in the shape of welding electrodes with a
diameter of 3.25 mm. The welding involved using a
direct electric current (+) with 110 A and 22 V. The
welding with preheating was performed only when
welding with the consumable EVB CrMo.
3.2 Microstructural characterization
The microstructural characterization of the base and
filler materials (as-welded and heat-treated state) was
performed using light optical microscopy (LOM –
ZEISS Axio Imager. A1m) and scanning electron micro-
scopy (SEM – JEOL JSM 5610). The emphasis was on
the characterization of the specimens on the side where
the fracture occurred. The metallographic specimens
were etched using a solution of 2 % nital. Prior to the
etching all three specimens were polished. The
microstructural characterization of the tensile-test
specimens was also made using the SEM on the fracture
surfaces in the welds and on the spots of the welds close
to the fracture.
3.3 Hardness measurements
The hardness measurements for each consumable
were made on a Shimadzu Micro Hardness Tester with a
load of 25 g and an indentation time of 15 s. The measu-
rements were made on both basic materials, both
heat-affected zones and the filling material on all three
specimens. The measurements were performed on
specimens in the as-welded condition and specimens in
the heat-treated state.
3.4 Tensile Test
The tensile tests were made on an INSTRON 1255.
For these tests the standard specimens were used. No
tensile tests were performed on the heat-treated speci-
mens.
3.5 Heat treatment
After the metallographic analysis the specimens were
heat treated in a chamber electrical resistance furnace.
The specimens were heated to 600 °C for one hour and
then cooled in the air. After the heat treatment the
metallographic examinations and the measurements of
the hardness were made on these specimens as well.
4 RESULTS AND DISCUSSION
4.1 Microstructural characterization
In the case of the consumable EVB 50 in the region
of the weld the microstructure consisted of ferrite and
bainite and in heat-affected zone (HAZ) there was also a
ferritic and bainitic microstructure. So, the weldment
consists mainly of ferrite and bainite (Figure 2a). The
microstructure of the base material EN 10216-1 mainly
consisted of ferrite. The heat treatment did not signifi-
cantly affect the microstructure in any part of the weld-
N. [TREKELJ ET AL.: METALLOGRAPHIC CHARACTERIZATION OF THE JOINING PART ...
408 Materiali in tehnologije / Materials and technology 44 (2010) 6, 407–412
Table 1: Chemical composition of steel EN 10216-1 and EN 10083-3 (mass fraction, w/%)1
Tabela 1: Kemi~na sestava jekla EN 10216-1 in EN 10083-3 (mas. dele`i, w/%)1
Material/ Element C Si Mn P S Cr Mo
EN 10216-1 0.115 0.270 1.119 / / / /
EN 10083-3 0.40 − 0.45 up to 0.40 0.6 − 0.9 up to 0.025 up to 0.035 0.9 − 1.2 0.15 − 0.3
Table 2: Chemical composition of consumables (mass fractions, w/%)5
Tabela 2: Kemi~na sestava dodajnih materialov (mas. dele`i, w/%)5
Consumable/Element C Si Mn Cr Mo Ni
EVB 50 0.07 0.60 1.00 / / /
EVB CrMo 0.06 0.60 0.95 1.10 0.50 /
INOX 29/9 0.11 up to 0.90 0.90 29.00 / 9.00
Figure 1: The section through the weldment1
Slika 1: Videz zvara po varjenju1
ment or the HAZ (Figure 2b). Figure 3a shows the
region of the weld with the consumable EVB CrMo,
which consists of ferrite and bainite. The HAZ shows a
ferritic microstructure with a small amount of bainite.
The heat treatment affected the columnar dendrite-like
morphology7 (Figure 3b). The weld with the consu-
mable INOX R 29/9 in the as-welded state is shown in
Figure 4a and in the heat-treated state in Figure 4b.
According to the Schäffler diagram,2 in the micro-
structure there should be austenite and ferrite. Figure 4a
shows the HAZ, which consists of ferrite, and the weld,
which consists of ferrite and pearlite. The microstructure
shows a dendritic morphology and the fusion line is
much thinner than in the previous two cases because of
the different compositions of the consumable and the
base materials. The diffusion of chromium and nickel
from the consumable INOX R 29/9 is limited to the area
of the fusion line, which was confirmed by the EDS spot
analyses and points to microsegregation. After the heat
treatment no changes were observed. In Figure 4b the
line between the weld and the base material EN 10216-1
probably shows the diffusion of the carbon to the
interface. The weldment is free of typical welding
defects, such as inclusions, porosity, microcracks etc., in
all three cases.
4.2 Hardness
Consumable EVB 50
The hardness measurements are in Figure 5 in
conjunction with the mean hardness, calculated in Table
4a. The hardness measurements (Figure 5) showed that
the average hardness of the base material EN 10083-3 in
the as-welded state was 	270 HV and for the EN
10216-1 it was 	180 HV. In the heat-treated state the
average hardness of the base material EN 10083-3 was
	 260 HV and for the base material EN 10216-1 it was
	240 HV. The base material EN 10216-1 and consu-
mable EVB 50 had similar compositions; therefore, they
had similar hardness values. The maximum hardness was
in the heat-affected zone (HAZ) on the side of the base
material EN 10083-3 (	420 HV). It slowly increased
near the HAZ (on the side of EN 10083-3), then sharply
increased near the fusion line, where it reached a
maximum of 	420 HV. Through the fusion line the
hardness decreased and reached an average hardness of
	200 HV in the weld. In the HAZ on the side of the EN
N. [TREKELJ ET AL.: METALLOGRAPHIC CHARACTERIZATION OF THE JOINING PART ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 407–412 409
Figure 4: Microstructure of consumable INOX R 29/9 on the side of
the base material EN 10216-1 in as-welded state (a) LM (left figure)
and SEM (right figure) and in heat-treated state (b) LM (left figure)
and SEM (right figure)
Slika 4: Mikrostruktura dodajnega materiala INOX R 29/9 na strani
osnovnega materiala EN 10216-1 v osnovnem stanju (a) SM (leva
slika) in SEM (desna slika) in v toplotno obdelanem stanju (b) SM
(leva slika) in SEM (desna slika)
Figure 3: Microstructure of consumable EVB CrMo on the side of the
base material EN 10216-1 in as-welded state (a) LM (left figure) and
SEM (right figure) and in heat-treated state (b) LM (left figure) and
SEM (right figure)
Slika 3: Mikrostruktura dodajnega materiala EVB CrMo (desno) na
strani osnovnega materiala EN 10216-1 (levo) v osnovnem stanju –
SM in SEM (a) in v toplotno obdelanem stanju – SM in SEM (b)
Figure 2: Microstructure of consumable EVB 50 on the side of the
base material EN 10216-1 as-welded state (a) LM (left figure) and
SEM (right figure) and in heat-treated state (b) LM (left figure) and
SEM (right figure)
Slika 2: Mikrostruktura dodajnega materiala EVB 50 na strani
osnovnega materiala EN 10216-1 v osnovnem stanju (a) SM (leva
slika) in SEM (desna slika) in v toplotno obdelanem stanju (b) SM
(leva slika) in SEM (desna slika)
10216-1 the hardness slightly decreased, where it was
around 150–200 HV.
In the heat-treated state the overall hardness was
lower on the side of the base material EN 10083-3.
Nevertheless, the hardness of the weld itself was higher
with 	280 HV. On the side of the base material EN
10216-1 the overall hardness was higher in heat-treated
state than in the as-welded state. The hardness of the
weld was also the highest hardness measured in the
heat-treated state of the weldment with the consumable
EVB 50.
Table 3: Ultimate tensile strengths
Tabela 3: Natezne trdnosti ob poru{itvi
Side of the base
material Consumable
Ultimate strength
/MPa
EN 10216-1 EVB 50 520
EN 10216-1 EVB CrMo 590
EN 10216-1 INOX R 29/9 550
Table 4: Mean hardness measurements of the weldments: a) EVB 50;
b) EVB CrMo anc c) INOX R 29/9
Tabela 4: Povpre~ne vrednosti trdot zvarjencev: a) EVB 50; b) EVB
CrMo in c) INOX R 29/9
a)
EVB 50 EN10083-3 HAZ Weld HAZ
EN
10216-1
As-welded 270 340 200 180 180
Heat-treated 260 270 280 210 240
b)
EVB CrMo EN10083-3 HAZ Weld HAZ
EN
10216-1
As-welded 280 360 260 260 200
Heat-treated 260 270 280 270 250
c)
INOX R 29/9 EN10083-3 HAZ Weld HAZ
EN
10216-1
As-welded 260 330 270 200 160
Heat-treated 270 320 360 230 260
Consumable EVB CrMo
The hardness measurements (Figure 6 and Table 4b)
in the as-welded state showed that the highest hardness
was in the HAZ on the side of the base material EN
10083-3 (	460 HV). The hardness of the base material
EN 10216-1 in as-welded state was 	200 HV, in
heat-treated state it was 	250 HV, and in the base
material EN 10083-3 it was 	280 HV in as-welded
state, and in the heat treated state it was 	260 HV. On
the transition from weld to the base material EN 10083-3
the hardness sharply increased to 400 HV and then it
decreased through the HAZ to 	260 HV in the weld.
In the heat-treated state the hardness was generally
decreased on the side of the base material EN 10083-3.
In the heat-treated state the highest average hardness was
in the weld (	280 HV). On the other side of the weld
(base material EN 10216-1) the hardness decreased. In
N. [TREKELJ ET AL.: METALLOGRAPHIC CHARACTERIZATION OF THE JOINING PART ...
410 Materiali in tehnologije / Materials and technology 44 (2010) 6, 407–412
Figure 7: Hardness of the weldment with the consumable INOX R
29/9 (blue line – as-welded state; red line – heat-treated state)
Slika 7: Trdote zvarjenca z dodajnim materilom INOX R 29/9 (modra
~rta – v osnovnem stanju; rde~a ~rta – v toplotno obdelanem stanju)
Figure 6: Hardness of the weldment with the consumable EVB CrMo
(blue line – as-welded state; red line – heat-treated state)
Slika 6: Trdote zvarjenca z dodajnim materialom EVB CrMo (modra
~rta – v osnovnem stanju; rde~a ~rta – v toplotno obdelanem stanju)
Figure 5: Hardness of the weldment with the consumable EVB 50
(blue line – as-welded state; red line – heat-treated state)
Slika 5: Trdote zvarjenca z dodajnim materialom EVB 50 (modra ~rta
– v osnovnem stanju; rde~a ~rta – v toplotno obdelanem stanju)
this case the hardness in the heat-treated state followed
the same trend as in the previous case and was higher on
the side of the base material EN 10216-1 than in
as-welded state.
Consumable INOX R 29/9
In this case the major difference was when the
specimen was heat treated (Figure 7 and Table 4c). The
hardness in the HAZ slightly increased (on the side of
the base material EN 10216-1) and in the weld it
strongly increased. This is due to the chromium (29 %)
in the consumable and the diffusion of the carbon into
the filler material. On the other side of the weld the
hardness again decreased slowly. The average hardness
of the base material EN 10216-1 was 	260 HV, and for
the base material EN 10083-3 it was 	270 HV.
In the as-welded state the hardness of the base
material EN 10083-3 was 	260 HV and of the base
material EN 10216-1 it was 	160 HV. In the weld the
hardness was 	270 HV, but the maximum average
hardness was in the HAZ (	330 HV) on the side of the
base material EN 10083-3. Again, the average hardness
in the heat-treated state was higher than in the as-welded
state. In all three cases the hardness in the heat-treated
state obviously increased due to the precipitation
hardening (on the side of the base material EN 10216-1).
4.3 Tensile Test
All the weldments were tested on an INSTRON
1255. The specimens were loaded until the fracture
occurred. After that the fracture surfaces were examined
using an SEM in order to establish the nature of the
fracture mechanics. In Figure 8 the fracture surfaces for
all three consumables are presented and in Table 3 the
ultimate tensile strengths are shown.
In all three cases the fracture occurred on the side
with the base material EN 10216-1. Therefore, the
characterization was focused on the sides of the fracture.
It is clear that in all three cases the fractures are ductile,
which is evident from Figure 8.
5 CONCLUSIONS
From the results of the tensile tests it is clear that the
problems occur on the side with the base material EN
10216-1, which is why the attention was focused on that
side of the weld. From the results of the microscopic
investigations of the welds with three base materials it
can be seen that the microstructure of the base material
EN 10216-1 consisted mainly of ferrite with some
pearlite (15%). The heat treatment did not significantly
affect the microstructure of the weldment with the
consumables EVB 50 and INOX R 29/9. However, the
heat treatment did change the columnar microstructure
of the weldment with the consumable EVB CrMo. The
EDS analysis of the weldment with INOX R 29/9
showed an increased concentration of chromium and
nickel close to the fusion line. This clearly showed the
presence of microsegregation. The hardness measure-
ments showed that the average hardness of the base
material EN 10083-3 in the as-welded state was 	275
HV and for the EN 10216-1 it was 	180 HV. The maxi-
mum hardness in the weldment with the consumable
EVB 50 was in the HAZ on the side of the base material
EN 10083-3 (	420 HV). The hardness measurements of
the weld with the consumable EVB CrMo in the
as-welded state showed that the highest hardness was in
the HAZ on the side of the base material EN 10083-3
(	460 HV). The hardness of the base material EN
10216-1 was 	200 HV. On the transition from the weld
to the base maerial EN 10083-3 the hardness sharply
increased to 400 HV and then it decreased through the
HAZ to 	260 HV in the weld. In the heat-treated state
N. [TREKELJ ET AL.: METALLOGRAPHIC CHARACTERIZATION OF THE JOINING PART ...
Materiali in tehnologije / Materials and technology 44 (2010) 6, 407–412 411
Figure 8: Fracture surfaces, examined using SEM; a) consumable
EVB 50, b) consumable EVB CrMo and c) consumable INOX R 29/9
Slika 8: Povr{ine prelomov, preiskovane s SEM; a) dodajni material
EVB 50, b) dodajni material EVB CrMo in c) dodajni material INOX
R 29/9
the highest average hardness was in the weld with 	280
HV. In the case with the consumable INOX R 29/9 the
major difference was when the specimen was heat
treated. The hardness in the weld strongly increased.
This was due to the chromium (29 %) in the consumable
and the diffusion of the carbon into the filler material. In
all three cases precipitation hardening was present
because the hardness increased in the heat-treated state
(in comparison with the as-welded state). The consu-
mable EVB CrMo shows a slight advantage over the
consumable EVB 50 and INOX R 29/9, in terms of the
ultimate tensile strength. When welding EN 10216-1 and
EN 10083-3 it is slightly better to weld with the con-
sumable EVB CrMo.
6 REFERENCES
1 M. Lampi~, Metallography of the joining part for the EN 10216-1
and EN 10083-3 made with different consumables, Diploma work (in
Slovene), Ljubljana 2009
2 R. Vrban, Metallography of the welded joint of an alloyed steel,
Diploma work (in Slovene), Ljubljana 2004
3 H. Schabereiter: Grundregeln für die Auswahl von Schweisszusatzen
für Mischverbindungen, Schweis und prüftechnik 10/98, Kapfenberg
4 W. Wuich, Nicht immer mit gleichen Schweisszusatz verwerden –
Probleme beim Schweissen von Gelenkwellen, 1996
5 Elektrode Jesenice: Consumables for welding, SIJ – Slovenian
industry of steel
6 V. Grdun, B. Godec, Unfavourable micro-constituents in the welded
joints of construction steel, Mater. Tehnol., 36 (2002) 5, 247–254
7 R. Celin, J. Tu{ek, D. Kmeti~, J. Vojvodi~ Tuma, An investigation of
welding of tool-steel, Mater. Tehnol., 35 (2001) 6, 405–408
8 S. Kastelic, P. Mrvar, J. Medved, M. Slova~ek: Simulation of the
welding and heat treatment process for S355 steel, Livarski vestnik,
56 (2009) 1, 25–38
N. [TREKELJ ET AL.: METALLOGRAPHIC CHARACTERIZATION OF THE JOINING PART ...
412 Materiali in tehnologije / Materials and technology 44 (2010) 6, 407–412
LETNO KAZALO – INDEX
Letnik / Volume 44 2010
ISSN 1580-2949
© Materiali in tehnologije
IMT Ljubljana, Lepi pot 11, 1000 Ljubljana, Slovenija
M
EHNOLOGIJEIN
AT E R IALI
M A T E R I A L S A N D T E C H N O L O G Y
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY
VSEBINA / CONTENTS
LETNIK / VOLUME 44, 2010/1, 2, 3, 4, 5, 6
2010/1
The thermal conductivity of Al73Mn27–xFex Taylor phases
Toplotna prevodnost Taylorjevih faz Al73Mn27–xFex
D. Stani}, P. Pop~evi}, I. Smiljani}, @. Bihar, J. Lukatela, B. Leonti}, A. Bilu{i}, I. Batisti}, A. Smontara . . . . . . . . . . . . . . . . . . . . . . . 3
Hall effect in the crystalline orthorombic o-Al13Co4 approximant to the decagonal quasicrystals
Hallov efekt v kristalini~nem ortorombi~nem pribli`ku o-Al13Co4 dekagonalnim kvazikristalom
J. Ivkov, D. Stani}, P. Pop~evi}, A. Smontara, J. Dolin{ek, P. Gille . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
Experimental plans method to formulate a self-compacting cement paste
Eksperimentalno na~rtovana metoda za dolo~itev samozgo{~ujo~e cementne malte
A. Mebrouki, N. Belas, N. Bouhamou . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13
The corrosion behaviour of austenitic and duplex stainless steels in artificial body fluids
Korozijsko vedenje avstenitnega in dupleksnega nerjavnega jekla v simuliranih telesnih teko~inah
A. Kocijan, M. Conradi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21
Chronic stimulation of an autonomous nerve with platinum electrodes
Kroni~na stimulacija avtonomnega `ivca s platinastimi elektrodami
P. Pe~lin, M. Zdovc, J. Rozman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25
Preparation of Si3N4-TiN ceramic composites
Priprava kerami~nih kompozitov na osnovi Si3N4-TiN
A. Maglica, K. Krnel, T. Kosma~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31
Analysis of the material and the actuator influence on the characteristics of a pneumatic valve
Analiza vpliva materiala in aktuatorjev na lastnosti pnevmati~nega ventila
N. Herakovi~, T. Bevk. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37
Stress analysis of a unilateral complex partial denture using the finite-element method
Napetostna analiza unilateralno kompleksne zobne proteze z uporabo metode kon~nih elementov
A. Todorovi}, K. Radovi}, A. Grbovi}, R. Rudolf, I. Maksimovi}, D. Stamenkovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 41
2010/2
Nekaj mo`nih postopkov izdelave hidrofobnih in oleofobnih polimernih membram ter njihova uporaba
Hydrophobic and olephobic membrane usage and production processes
D. Drev, J. Panjan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51
Liquid metal/ceramic interfaces in dental practice and jewellery manufacturing
Staljene vmesne povr{ine med kovino in keramiko v zobozdravstvu in pri izdelavi nakita
K. T. Rai}, R. Rudolf, A. Todorovi}, D. Stamenkovi}, I. An`el . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59
Influence of MnS inclusions on the corrosion of austenitic stainless steels
Vpliv vklju~kov MnS na korozijsko odpornost avstenitnih nerjavnih jekel
^. Donik, I. Paulin, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67
AES and XPS characterization of titanium hydride powder
Preiskave pra{ka titanovega hidrida s spektroskopijo Augerjevih elektronov in rentgensko fotoelektronsko spektroskopijo
I. Paulin, D. Mandrino, ^. Donik, M. Jenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73
Fracture characteristics of the Cr-V ledeburitic steel Vanadis 6
Prelomne zna~ilnosti ledeburitnega Cr-V-jekla Vanadis 6
P. Jur~i, B. [u{tar{i~, V. Leskov{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77
Surface modifications of maraging steels used in the manufacture of moulds and dies
Modifikacija povr{ine jekla maraging in uporaba pri izdelavi kokil in utopov
F. Cajner, D. Landek, V. Leskov{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85
Two numerical models of the solidification structure of massive ductile cast-iron casting
Numeri~na modela strukture strjevanja masivnega duktilnega `elezovega ulitka
K. Stransky, J. Dobrovska, F. Kavicka, V. Gontarev, B. Sekanina, J. Stetina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93
414 Materiali in tehnologije / Materials and technology 44 (2010) 6
LETNO KAZALO – INDEX
Correlation between the corrosion resistance and the hardness scattering of structural metals treated with a pulsed electric
current
Korelacija med odpornostjo proti koroziji in raztrosom trdote pri konstrukcijskih materialih, obdelanih s pulziranjem elektri~nega toka
A. Babutsky, A. Chrysanthou, J. Ioannou, I. Mamuzic . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99
2010/3
Physical simulation of metallurgical processes
Fizikalna simulacija metalur{kih procesov
S. T. Mandziej . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 105
Thermodynamic modeling for the alloy design of high speed steels and high chromium cast irons
Termodinami~no modeliranje na~rtovanja sestave hitroreznih jekel in litine z veliko vsebnostjo kroma
M. Pellizzari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 121
Use of Grey based Taguchi method in ball burnishing process for the optimization of surface roughness and microhardness of
AA 7075 aluminum alloy
Uporaba Grey-Taguchijeve metode pri procesu glajenja za optimizacijo povr{inske hrapavosti in mikrotrdote aluminijeve zlitine AA 7075
U. Esme . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 129
Effect of cerium additions on the AlSi17 casting alloy
Dodatek cerija livni zlitini AlSi17
S. Kores, M. Von~ina, B. Kosec, P. Mrvar, J. Medved . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 137
A generalized theory of plasticity
Posplo{ena teorija plasti~nosti
V. V. Chygyryns’kyy, K. A. Yakovlevich, I. Mamuzi}, B. A. Nikolaevna . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 141
Comparative modeling of wire electrical discharge machining (WEDM) process using back propagation (BPN) and general
regression neural networks (GRNN)
Primerjalno modeliranje elektroerozijske `i~ne obdelave (WEDM) z uporabo povratnosti (BPN) in spo{ne nevronske regresijske mre`e
(GRNN)
O. Guven, U. Esme, I. E. Kaya, Y. Kazancoglu, M. K. Kulekci, C. Boga . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147
Degradation of bacteria Escherichia coli by treatment with Ar ion beam and neutral oxygen atoms
Uni^evanje bakterij Escherichia coli s curkom ionov Ar in nevtralnih atomov kisika
K. Eler{i~, I. Junkar, A. [pes, N. Hauptman, M. Klanj{ek-Gunde, A. Vesel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153
Effect of heat treatment on the microstructure and properties of Cr-V ledeburitic steel
Vpliv toplotne obdelave na mikrostrukturo in lastnosti ledeburitnega jekla Cr-V
S. Krum, J. Sobotová, P. Jur~i . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157
2010/4
Surface modification of materials using an extremely non-equilibrium oxygen plasma
Modifikacija povr{ine materialov z izrazito neravnovesno kisikovo plazmo
M. Mozeti~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 165
Middle triassic dry-land phases in Southern Slovenia
Srednjetriasne kopenske faze v ju`ni Sloveniji
S. Dozet, M. Godec. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 173
Enhanced fatigue analysis – incorporating downstream manufacturing processes
Izbolj{ana utrujenostna analiza – vklju~itev izdelavnih procesov
W. Eichlseder . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 185
A new complex vector method for balancing chemical equations
Nova kompleksna metoda za uravnote`enje kemi~nih ena~b
I. B. Risteski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 193
Application of Grey relation analysis (GRA) and Taguchi method for the parametric optimization of friction stir welding
(FSW) process
Uporaba Greyjeve analize (GRA) in Taguchijeve metode za parametri~no optimizacijo varjenja z vrtilno-tornim procesom (FSW)
H. Aydin, A. Bayram, U. Esme, Y. Kazancoglu, O. Guven . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 205
Phase relations in the Pt–Ag17.5–Si4.5 ternary alloy
Fazna odvisnost v ternarni zlitini Pt–Ag17.5–Si4.5
G. Klan~nik, P. Mrvar, J. Medved . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 213
A new solution of the harmonic functions in the theory of elasticity
Nekatere zna~ilnosti harmoni~nih funkcij v teoriji o elastni~nosti
V. V. Chygyryns’kyy, V. G. Shevchenko, I. Mamuzic, S. B. Belikov. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 219
Materiali in tehnologije / Materials and technology 44 (2010) 6 415
LETNO KAZALO – INDEX
Influence of vacuum processing on the content of some elements in molten metal
Vpliv vakuumskega procesiranja na vsebnost nekaterih elementov v kovinskih talinah
D. Pihura, D. Mujagi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 223
Influence of -ferrite on the fatigue resistance of blade materials
Vpliv -ferita na utrujenostno trdnost jekel za lopatice
P. Hájková, J. Janovec, J. Siegl, B. Smola, M. Vyroubalová, D. Tùmová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227
2010/5
Selective oxidation of manganese in molten pig iron
Selektivna oksidacija mangana v staljenem grodlju
D. Pihura, M. Oru~, J. Lamut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 235
The corrosion behaviour of ferritic stainless steels in alkaline solutions
Korozijsko vedenje feritnih nerjavnih jekel v alkalnih raztopinah
A. Kocijan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 239
Numerical design of a hot-stretch-reducing process for welded tubes
Numeri~no na~rtovanje vro~e raztezne redukcije varjenih cevi
S. Re{kovi}, R. Kri`ani}, F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243
Preparation and study of Mg2Sn-based composites with different compositions
Priprava in karakterizacija lastnosti kompozitov na osnovi Mg2Sn
V. Kevorkijan, S. D. [kapin. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 251
Osseointegration and rejection of a titanium screw
Oseointegracija in zavrnitev titanovega vijaka
G. Klan~nik, M. Zdovc, U. Kov{ca, B. Pra~ek, J. Kova~, J. Rozman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 261
Corrosion investigation of the Pre{eren monument in Ljubljana
Korozijska preiskava Pre{ernovega spomenika v Ljubljani
M. Bajt Leban, V. Kuhar, T. Kosec, A. Legat . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 265
The microflora and physico-chemical properties of lignite from the Mirash mine, near Kastriot
Mikroflora in fizikalno-kemijske lastnosti lignita iz rudnika Mirash pri Kastriotu
F. Plakolli, L. P. Beqiri, A. Millaku . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 271
Cadmium and lead accumulation in Alfalfa (Medicago Sativa L.) and their influence on the number of stomata
Akumulacija kadmija in svinca v rastlini Lucerna (Medicago Sativa L.) in njun vpliv na {tevilo listnih por
A. Bytyqi, E. Sherifi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 277
The selection and development of tribological coatings
Izbira in razvoj tribolo{kih prevlek
Y. Kharlamov, V. Dal, I. Mamuzi}, L. Lopata, G. S. Pisarenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283
Recycling of steel chips
Recikliranje jeklenih ostru`kov
M. Torkar, M. Lamut, A. Millaku . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289
2010/6
In memoriam – Jo`ef Arh 1930–2010 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295
60 let In{tituta za kovinske materiale in tehnologije Ljubljana
60 years of the Institute of metals and technology Ljubljana
M. Jenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 297
40 let elektronske mikroanalize
40 years of electron probe mikroanalisis
F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 305
25 let vakuumske toplotne in kemotermi~ne obdelave na In{titutu za kovinske materiale in tehnologije, Ljubljana
25 years of vacuum heat treatment and surface engineering at the Institute of metals and technology, Ljubljana
V. Leskov{ek. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 307
Non-oriented electrical steel sheets
Neorientirane elektroplo~evine
D. Steiner Petrovi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 317
Effect of trace and residual elements on the hot brittleness, hot shortness and properties of 0.15–0.3 % C Al-killed steels with
a solidification microstructure
U~inek elementov v sledovih in preostalih elementov na krhkost in pokljivost v vro~em jekla z 0,15 – 0,3 % C, pomirjenega z Al in s
strjevalno mikrostrukturo
M. Torkar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 327
416 Materiali in tehnologije / Materials and technology 44 (2010) 6
LETNO KAZALO – INDEX
The influence of powder selection and process conditions on the characteristics of some oil-retaining sintered bronze bearings
Vpliv izbire prahov in pogojev izdelave na lastnosti samomazalnih sintranih le`ajev vrste Cu-Sn
B. [u{tar{i~, M. Godec, A. Kocijan, ^. Donik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 335
The influence of the microalloying elements of HSLA steel on the microstructure and mechanical properties
Vpliv mikrolegirnih elementov na mikrostrukturo in mehanske lastnosti HSLA jekla
D. A. Skobir, M. Godec, M. Balcar, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 343
Effect of 3 S reheating on charpy notch toughness and hardness of martensite and lower bainite
Vpliv 3 S pogrevanja na charpyjevo zarezno `ilavost in trdoto martenzita in spodnjega bainita
G. Kosec, F. Vodopivec, A. Smolej, J. Vojvodi~-Tuma, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 349
Deformation mechanisms in Ti3Al-Nb alloy at elevated temperatures
Mehanizem deformacije v zlitini Ti3Al-Nb pri povi{anih temperaturah
S. Tadi}, R. Proki}-Cvetkovi}, I. Bala}, R. Heinemann-Jan~i}, K. Boji}, A. Sedmak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 357
Synthesis and characterisation of closed cells aluminium foams containing dolomite powder as foaming agent
Priprava in karakterizacija aluminijskih pen z zaprto poroznostjo, izdelanih z dolomitnim prahom kot sredstvom za penjenje
V. Kevorkijan, S. Davor [kapin, I. Paulin, B. [u{tar{i~, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 363
Imobilizacija lete~ega pepela z zasteklitvijo
Fly ash immobilization with vitrification
N. Ore{ek, F. Berk, N. Samec, F. Zupani~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 373
The impact of stagnant water on the corrosion processes in a pipeline
Vpliv zastajajo~ih voda v cevovodih na korozijske procese
M. Suban, R. Cvelbar, B. Bundara. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 379
Determination of the critical pressure for a hot-water pipe with a corrosion defect
Dolo~itev kriti~nega pritiska v vro~evodni cevi s korozijsko po{kodbo
D. Kozak, @. Ivandi}, P. Kontaji}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 385
Effects of variations in alloy content and machining parameters on the strength of the intermetallic bonding between a diesel
piston and a ring carrier
Vpliv sprememb v sestavi zlitine in parametrov obdelave na trdnost intermetalne vezave med dieselskim batom in nosilcem obro~ka
A. F. Acar, F. Ozturk, M. Bayrak. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 391
Analysis of the process of crystallization of continuous cast special brass alloys with the acoustic emission
Analiza procesa kristalizacije kontinuirno lite specialne medenine z metodo akusti~ne emisije
M. Tasi}, Z. An|i}, A. Vujovi}, P. Jovani}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 397
The influence of adding emulsion flocculants on the effects of red-mud sedimentation
Vpliv dodatka emulzijskih flokulantov na pojave pri sedimentaciji rde~ega blata
D. Smolovi}, M. Vuk~evi}, D. Ble~i} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 403
Metallographic characterization of the joining part for EN 10216-1 and EN 10083-3 made with different consumables
Metalografska karakterizacija veznega ~lena iz EN 10216-1 in EN 10083-3, izdelanega z razli~nimi dodajnimi materiali
N. [trekelj, M. Lampi~, L. Kosec, B. Markoli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 407
Materiali in tehnologije / Materials and technology 44 (2010) 6 417
LETNO KAZALO – INDEX
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY
AVTORSKO KAZALO / AUTHOR INDEX
LETNIK / VOLUME 44, 2010, 1–6, A–@
A
An`el I. 59
An|i} Z. 397
Aydin H. 205
B
Babutsky A. 99
Bajt Leban M. 265
Bala} I. 357
Balcar M. 343
Batisti} I. 3
Bayrak M. 391
Bayram A. 205
Belas N. 13
Belikov S. B. 219
Beqiri L. P. 271
Berk F. 373
Bevk T. 37
Bihar @. 3
Bilu{i} A. 3
Ble~i} D. 403
Boga C. 147
Boji} K. 357
Bouhamou N. 13
Bundara B. 379
Bytyqi A. 277
C
Cajner F. 85
Chrysanthou A. 99
Chygyryns’kyy V. V. 141, 219
Conradi M. 21
Cvelbar R. 379
D
Dal V. 283
Davor [kapin S. 363
Dobrovska J. 93
Dolin{ek J. 9
Donik ^. 67, 73, 335
Dozet S. 173
Drev D. 51
E
Eichlseder W. 185
Eler{i~ K. 153
Esme U. 129, 147, 205
Fahriye Acar A. 391
G
Gille P. 9
Godec M. 173, 335, 343
Gontarev V. 93
Grbovi} A. 41
Guven O. 147, 205
H
Hájková P. 227
Hauptman N. 153
Heinemann-Jan~i} R. 357
Herakovi~ N. 37
I
Ioannou J. 99
Ivandi} @. 385
Ivkov J. 9
J
Janovec J. 227
Jenko M. 67, 73, 297, 343, 349, 363
Jovani} P. 397
Junkar I. 153
Jur~i P. 77, 157
K
Kavicka F. 93
Kaya I. E. 147
Kazancoglu Y. 147, 205
Kevorkijan V. 251, 363
Kharlamov Y. 283
Klan~nik G. 213, 261
Klanj{ek-Gunde M. 153
Kocijan A. 21, 239, 335
Kontaji} P. 385
Kores S. 137
Kosec B. 137
Kosec G. 349
Kosec L. 407
Kosec T. 265
Kosma~ T. 31
Kov{ca U. 261
Kova~ J. 261
Kozak D. 385
Kri`ani} R. 243
Krnel K. 31
Krum S. 157
Kuhar V. 265
Kulekci M. K. 147
L
Lampi~ M. 407
Lamut J. 235
Lamut M. 289
Landek D. 85
Legat A. 265
Leonti} B. 3
Leskov{ek V. 77, 85, 307
Lopata L. 283
Lukatela J. 3
M
Maglica A. 31
Maksimovi} I. 41
Mamuzic I. 99, 141, 283, 219
Mandrino D. 73
Mandziej S. T. 105
Markoli B. 407
Mebrouki A. 13
Medved J. 137, 213
Millaku A. 271, 289
Mozeti~ M. 165
Mrvar P. 137, 213
Mujagi} D. 223
N
Nikolaevna B. A. 141
O
Ore{ek N. 373
Oru~ M. 235
Ozturk F. 391
P
Panjan J. 51
Paulin I. 67, 73, 363
Pe~lin P. 25
Pellizzari M. 121
Pihura D. 223, 235
Pisarenko G. S. 283
Plakolli F. 271
418 Materiali in tehnologije / Materials and technology 44 (2010) 6
LETNO KAZALO – INDEX
Pop~evi} P. 3, 9
Pra~ek B. 261
Proki}-Cvetkovi} R. 357
R
Radovi} K. 41
Rai} K. T. 59
Re{kovi} S. 243
Risteski I. B. 193
Rozman J. 25, 261
Rudolf R. 41, 59
S
Samec N. 373
Sedmak A. 357
Sekanina B. 93
Sherifi E. 277
Shevchenko V. G. 219
Siegl J. 227
Skobir D. A. 343
Smiljani} I. 3
Smola B. 227
Smolej A. 349
Smolovi} D. 403
Smontara A. 3, 9
Sobotová J. 157
Stamenkovi} D. 41, 59
Stani} D. 3, 9
Steiner Petrovi~ D. 317
Stetina J. 93
Stransky K. 93
Suban M. 379
S
[kapin S. D. 251
[pes A. 153
[trekelj N. 407
[u{tar{i~ B. 77, 335, 363
T
Tadi} S. 357
Tasi} M. 397
Todorovi} A. 41, 59
Torkar M. 289, 327
Tumová D. 227
V
Vesel A. 153
Vodopivec F. 243, 305, 349
Vojvodi~-Tuma J. 349
Von~ina M. 137
Vujovi} A. 397
Vuk~evi} M. 403
Vyroubalová M. 227
Y
Yakovlevich K. A. 141
Z
Zdovc M. 25, 261
Zupani~ F. 373
Materiali in tehnologije / Materials and technology 44 (2010) 6 419
LETNO KAZALO – INDEX
MATERIALI IN TEHNOLOGIJE / MATERIALS AND
TECHNOLOGY
VSEBINSKO KAZALO / SUBJECT INDEX
LETNIK / VOLUME 43, 2009, 1–6
Kovinski materiali – Metallic materials
The thermal conductivity of Al73Mn27–xFex Taylor phases
Toplotna prevodnost Taylorjevih faz Al73Mn27–xFex
D. Stani}, P. Pop~evi}, I. Smiljani}, @. Bihar, J. Lukatela, B. Leonti}, A. Bilu{i}, I. Batisti}, A. Smontara . . . . . . . . . . . . . . . . . . . . . . . 3
Hall effect in the crystalline orthorombic o-Al13Co4 approximant to the decagonal quasicrystals
Hallov efekt v kristalini~nem ortorombi~nem pribli`ku o-Al13Co4 dekagonalnim kvazikristalom
J. Ivkov, D. Stani}, P. Pop~evi}, A. Smontara, J. Dolin{ek, P. Gille . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
The corrosion behaviour of austenitic and duplex stainless steels in artificial body fluids
Korozijsko vedenje avstenitnega in dupleksnega nerjavnega jekla v simuliranih telesnih teko~inah
A. Kocijan, M. Conradi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21
Chronic stimulation of an autonomous nerve with platinum electrodes
Kroni~na stimulacija avtonomnega `ivca s platinastimi elektrodami
P. Pe~lin, M. Zdovc, J. Rozman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25
Analysis of the material and the actuator influence on the characteristics of a pneumatic valve
Analiza vpliva materiala in aktuatorjev na lastnosti pnevmati~nega ventila
N. Herakovi~, T. Bevk. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37
Stress analysis of a unilateral complex partial denture using the finite-element method
Napetostna analiza unilateralno kompleksne zobne proteze z uporabo metode kon~nih elementov
A. Todorovi}, K. Radovi}, A. Grbovi}, R. Rudolf, I. Maksimovi}, D. Stamenkovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 41
Liquid metal/ceramic interfaces in dental practice and jewellery manufacturing
Staljene vmesne povr{ine med kovino in keramiko v zobozdravstvu in pri izdelavi nakita
K. T. Rai}, R. Rudolf, A. Todorovi}, D. Stamenkovi}, I. An`el . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59
Influence of MnS inclusions on the corrosion of austenitic stainless steels
Vpliv vklju~kov MnS na korozijsko odpornost avstenitnih nerjavnih jekel
^. Donik, I. Paulin, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67
AES and XPS characterization of titanium hydride powder
Preiskave pra{ka titanovega hidrida s spektroskopijo Augerjevih elektronov in rentgensko fotoelektronsko spektroskopijo
I. Paulin, D. Mandrino, ^. Donik, M. Jenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73
Fracture characteristics of the Cr-V ledeburitic steel Vanadis 6
Prelomne zna~ilnosti ledeburitnega Cr-V-jekla Vanadis 6
P. Jur~i, B. [u{tar{i~, V. Leskov{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77
Surface modifications of maraging steels used in the manufacture of moulds and dies
Modifikacija povr{ine jekla maraging in uporaba pri izdelavi kokil in utopov
F. Cajner, D. Landek, V. Leskov{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85
Two numerical models of the solidification structure of massive ductile cast-iron casting
Numeri~na modela strukture strjevanja masivnega duktilnega `elezovega ulitka
K. Stransky, J. Dobrovska, F. Kavicka, V. Gontarev, B. Sekanina, J. Stetina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93
Correlation between the corrosion resistance and the hardness scattering of structural metals treated with a pulsed electric
current
Korelacija med odpornostjo proti koroziji in raztrosom trdote pri konstrukcijskih materialih, obdelanih s pulziranjem elektri~nega toka
A. Babutsky, A. Chrysanthou, J. Ioannou, I. Mamuzic . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99
Physical simulation of metallurgical processes
Fizikalna simulacija metalur{kih procesov
S. T. Mandziej . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 105
Thermodynamic modeling for the alloy design of high speed steels and high chromium cast irons
Termodinami~no modeliranje na~rtovanja sestave hitroreznih jekel in litine z veliko vsebnostjo kroma
M. Pellizzari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 121
420 Materiali in tehnologije / Materials and technology 44 (2010) 6
LETNO KAZALO – INDEX
Use of Grey based Taguchi method in ball burnishing process for the optimization of surface roughness and microhardness of
AA 7075 aluminum alloy
Uporaba Grey-Taguchijeve metode pri procesu glajenja za optimizacijo povr{inske hrapavosti in mikrotrdote aluminijeve zlitine AA 7075
U. Esme . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 129
Effect of cerium additions on the AlSi17 casting alloy
Dodatek cerija livni zlitini AlSi17
S. Kores, M. Von~ina, B. Kosec, P. Mrvar, J. Medved . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 137
A generalized theory of plasticity
Posplo{ena teorija plasti~nosti
V. V. Chygyryns’kyy, K. A. Yakovlevich, I. Mamuzi}, B. A. Nikolaevna . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 141
Comparative modeling of wire electrical discharge machining (WEDM) process using back propagation (BPN) and general
regression neural networks (GRNN)
Primerjalno modeliranje elektroerozijske `i~ne obdelave (WEDM) z uporabo povratnosti (BPN) in spo{ne nevronske regresijske mre`e
(GRNN)
O. Guven, U. Esme, I. E. Kaya, Y. Kazancoglu, M. K. Kulekci, C. Boga . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147
Effect of heat treatment on the microstructure and properties of Cr-V ledeburitic steel
Vpliv toplotne obdelave na mikrostrukturo in lastnosti ledeburitnega jekla Cr-V
S. Krum, J. Sobotová, P. Jur~i . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157
Surface modification of materials using an extremely non-equilibrium oxygen plasma
Modifikacija povr{ine materialov z izrazito neravnovesno kisikovo plazmo
M. Mozeti~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 165
Enhanced fatigue analysis – incorporating downstream manufacturing processes
Izbolj{ana utrujenostna analiza – vklju~itev izdelavnih procesov
W. Eichlseder . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 185
Application of Grey relation analysis (GRA) and Taguchi method for the parametric optimization of friction stir welding
(FSW) process
Uporaba Greyjeve analize (GRA) in Taguchijeve metode za parametri~no optimizacijo varjenja z vrtilno-tornim procesom (FSW)
H. Aydin, A. Bayram, U. Esme, Y. Kazancoglu, O. Guven . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 205
Phase relations in the Pt–Ag17.5–Si4.5 ternary alloy
Fazna odvisnost v ternarni zlitini Pt–Ag17.5–Si4.5
G. Klan~nik, P. Mrvar, J. Medved . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 213
A new solution of the harmonic functions in the theory of elasticity
Nekatere zna~ilnosti harmoni~nih funkcij v teoriji o elastni~nosti
V. V. Chygyryns’kyy, V. G. Shevchenko, I. Mamuzic, S. B. Belikov. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 219
Influence of vacuum processing on the content of some elements in molten metal
Vpliv vakuumskega procesiranja na vsebnost nekaterih elementov v kovinskih talinah
D. Pihura, D. Mujagi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 223
Influence of -ferrite on the fatigue resistance of blade materials
Vpliv -ferita na utrujenostno trdnost jekel za lopatice
P. Hájková, J. Janovec, J. Siegl, B. Smola, M. Vyroubalová, D. Tùmová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227
Selective oxidation of manganese in molten pig iron
Selektivna oksidacija mangana v staljenem grodlju
D. Pihura, M. Oru~, J. Lamut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 235
The corrosion behaviour of ferritic stainless steels in alkaline solutions
Korozijsko vedenje feritnih nerjavnih jekel v alkalnih raztopinah
A. Kocijan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 239
Numerical design of a hot-stretch-reducing process for welded tubes
Numeri~no na~rtovanje vro~e raztezne redukcije varjenih cevi
S. Re{kovi}, R. Kri`ani}, F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243
Preparation and study of Mg2Sn-based composites with different compositions
Priprava in karakterizacija lastnosti kompozitov na osnovi Mg2Sn
V. Kevorkijan, S. D. [kapin. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 251
Osseointegration and rejection of a titanium screw
Oseointegracija in zavrnitev titanovega vijaka
G. Klan~nik, M. Zdovc, U. Kov{ca, B. Pra~ek, J. Kova~, J. Rozman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 261
Corrosion investigation of the Pre{eren monument in Ljubljana
Korozijska preiskava Pre{ernovega spomenika v Ljubljani
M. Bajt Leban, V. Kuhar, T. Kosec, A. Legat . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 265
LETNO KAZALO – INDEX
Materiali in tehnologije / Materials and technology 44 (2010) 6 421
The selection and development of tribological coatings
Izbira in razvoj tribolo{kih prevlek
Y. Kharlamov, V. Dal, I. Mamuzi}, L. Lopata, G. S. Pisarenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283
Recycling of steel chips
Recikliranje jeklenih ostru`kov
M. Torkar, M. Lamut, A. Millaku . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289
25 let vakuumske toplotne in kemotermi~ne obdelave na In{titutu za kovinske materiale in tehnologije, Ljubljana
25 years of vacuum heat treatment and surface engineering at the Institute of metals and technology, Ljubljana
V. Leskov{ek. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 307
Non-oriented electrical steel sheets
Neorientirane elektroplo~evine
D. Steiner Petrovi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 317
Effect of trace and residual elements on the hot brittleness, hot shortness and properties of 0.15–0.3 % C Al-killed steels with
a solidification microstructure
U~inek elementov v sledovih in preostalih elementov na krhkost in pokljivost v vro~em jekla z 0,15 – 0,3 % C, pomirjenega z Al in s
strjevalno mikrostrukturo
M. Torkar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 327
The influence of powder selection and process conditions on the characteristics of some oil-retaining sintered bronze bearings
Vpliv izbire prahov in pogojev izdelave na lastnosti samomazalnih sintranih le`ajev vrste Cu-Sn
B. [u{tar{i~, M. Godec, A. Kocijan, ^. Donik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 335
The influence of the microalloying elements of HSLA steel on the microstructure and mechanical properties
Vpliv mikrolegirnih elementov na mikrostrukturo in mehanske lastnosti HSLA jekla
D. A. Skobir, M. Godec, M. Balcar, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 343
Effect of 3 S reheating on charpy notch toughness and hardness of martensite and lower bainite
Vpliv 3 S pogrevanja na charpyjevo zarezno `ilavost in trdoto martenzita in spodnjega bainita
G. Kosec, F. Vodopivec, A. Smolej, J. Vojvodi~-Tuma, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 349
Deformation mechanisms in Ti3Al-Nb alloy at elevated temperatures
Mehanizem deformacije v zlitini Ti3Al-Nb pri povi{anih temperaturah
S. Tadi}, R. Proki}-Cvetkovi}, I. Bala}, R. Heinemann-Jan~i}, K. Boji}, A. Sedmak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 357
Synthesis and characterisation of closed cells aluminium foams containing dolomite powder as foaming agent
Priprava in karakterizacija aluminijskih pen z zaprto poroznostjo, izdelanih z dolomitnim prahom kot sredstvom za penjenje
V. Kevorkijan, S. Davor [kapin, I. Paulin, B. [u{tar{i~, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 363
The impact of stagnant water on the corrosion processes in a pipeline
Vpliv zastajajo~ih voda v cevovodih na korozijske procese
M. Suban, R. Cvelbar, B. Bundara. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 379
Determination of the critical pressure for a hot-water pipe with a corrosion defect
Dolo~itev kriti~nega pritiska v vro~evodni cevi s korozijsko po{kodbo
D. Kozak, @. Ivandi}, P. Kontaji}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 385
Effects of variations in alloy content and machining parameters on the strength of the intermetallic bonding between a diesel
piston and a ring carrier
Vpliv sprememb v sestavi zlitine in parametrov obdelave na trdnost intermetalne vezave med dieselskim batom in nosilcem obro~ka
A. F. Acar, F. Ozturk, M. Bayrak. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 391
Analysis of the process of crystallization of continuous cast special brass alloys with the acoustic emission
Analiza procesa kristalizacije kontinuirno lite specialne medenine z metodo akusti~ne emisije
M. Tasi}, Z. An|i}, A. Vujovi}, P. Jovani}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 397
Metallographic characterization of the joining part for EN 10216-1 and EN 10083-3 made with different consumables
Metalografska karakterizacija veznega ~lena iz EN 10216-1 in EN 10083-3, izdelanega z razli~nimi dodajnimi materiali
N. [trekelj, M. Lampi~, L. Kosec, B. Markoli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 407
Anorganski materiali – Inorganic materials
Experimental plans method to formulate a self-compacting cement paste
Eksperimentalno na~rtovana metoda za dolo~itev samozgo{~ujo~e cementne malte
A. Mebrouki, N. Belas, N. Bouhamou . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13
Preparation of Si3N4-TiN ceramic composites
Priprava kerami~nih kompozitov na osnovi Si3N4-TiN
A. Maglica, K. Krnel, T. Kosma~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31
LETNO KAZALO – INDEX
422 Materiali in tehnologije / Materials and technology 44 (2010) 6
Middle triassic dry-land phases in Southern Slovenia
Srednjetriasne kopenske faze v ju`ni Sloveniji
S. Dozet, M. Godec. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 173
A new complex vector method for balancing chemical equations
Nova kompleksna metoda za uravnote`enje kemi~nih ena~b
I. B. Risteski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 193
The microflora and physico-chemical properties of lignite from the Mirash mine, near Kastriot
Mikroflora in fizikalno-kemijske lastnosti lignita iz rudnika Mirash pri Kastriotu
F. Plakolli, L. P. Beqiri, A. Millaku . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 271
Imobilizacija lete~ega pepela z zasteklitvijo
Fly ash immobilization with vitrification
N. Ore{ek, F. Berk, N. Samec, F. Zupani~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 373
The influence of adding emulsion flocculants on the effects of red-mud sedimentation
Vpliv dodatka emulzijskih flokulantov na pojave pri sedimentaciji rde~ega blata
D. Smolovi}, M. Vuk~evi}, D. Ble~i} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 403
Polimeri – Polymers
Nekaj mo`nih postopkov izdelave hidrofobnih in oleofobnih polimernih membram ter njihova uporaba
Hydrophobic and olephobic membrane usage and production processes
D. Drev, J. Panjan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51
Vakuumska tehnika – Vacuum technique
Degradation of bacteria Escherichia coli by treatment with Ar ion beam and neutral oxygen atoms
Uni^evanje bakterij Escherichia coli s curkom ionov Ar in nevtralnih atomov kisika
K. Eler{i~, I. Junkar, A. [pes, N. Hauptman, M. Klanj{ek-Gunde, A. Vesel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153
Varovanje okolja – Environmental protection
Cadmium and lead accumulation in Alfalfa (Medicago Sativa L.) and their influence on the number of stomata
Akumulacija kadmija in svinca v rastlini Lucerna (Medicago Sativa L.) in njun vpliv na {tevilo listnih por
A. Bytyqi, E. Sherifi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 277
Raziskovalni in{tituti – Research Institutes
60 let In{tituta za kovinske materiale in tehnologije Ljubljana
60 years of the Institute of metals and technology Ljubljana
M. Jenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 297
40 let elektronske mikroanalize
40 years of electron probe mikroanalisis
F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 305
Osebne biografije – Personal biographies
In memoriam – Jo`ef Arh 1930–2010 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295
LETNO KAZALO – INDEX
Materiali in tehnologije / Materials and technology 44 (2010) 6 423