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MINISTERUL EDUCAŢIEI ŞI CERCETĂRII

ANALELE UNIVERSITĂŢII “DUNĂREA DE JOS” DIN GALAŢI

Fascicula IX

FACULTATEA DE METALURGIE ŞI ŞTIINŢA MATERIALELOR

ANUL XXII (XXVII), mai. 2004, nr.1

ISSN 1453-083X

MINISTRY OF EDUCATION AND RESEARCH

THE ANNALS OF “DUNAREA DE JOS” UNIVERSITY OF GALATI

Fascicle IX

FACULTY OF METALLURGY AND MATERIALS SCIENCE

YEAR XXII (XXVII), May. 2004, no.1

ISSN 1453-083X

EDITING MANAGEMENT

RESPONSIBLE EDITOR: Prof. Dr. Eng. Alexandru EPUREANU

ASSISTANT EDITORS: Prof. Dr. Eng. Emil CONSTANTIN

Prof. Dr. Eng. Viorel MINZU Prof. Dr. Eng. Mircea BULANCEA Conf. Dr. Ec. Daniela ŞARPE Conf. Dr. Anca GÂŢĂ

SECRETARY: Assoc. Prof. Dr. Eng. Ion ALEXANDRU

EDITING BOARD

Fascicle IX

METALLURGY AND MATERIALS SCIENCE

EDITOR IN CHIEF: Prof. Dr. Chim. Olga Mitoşeriu

SECRETARY: Prof. Dr. Eng. Marian Bordei MEMBERS: Acad. Prof. Dr. Hab. Iurie Nicolaevich Shevcenko–Director of the Termoplasticity Department, National Academy of Science of Ukraine Acad. Prof. Dr. Hab. Valeriu Kantser–Coordinator of the Technical and Scientific Section of the Academy of Moldova Republic Prof. Dr. Rodrigo Martins – President of the Department of Materials Science, Faculty of Science and Technology, NOVA University of Lisbon, Portugal Prof.Dr.Hab. Vasile Marina–Director of Department, State Technical University of Moldova, Kishinau, Moldova Republic Prof. Dr. Eng. Elena Drugescu Prof. Dr. Eng. Nicolae Cănănău Prof. Dr. Eng. Anisoara Ciocan Prof. Dr. Eng. Maria Vlad Prof. Dr. Eng. Petre Stelian Niţă Prof. Dr. Eng. Alexandru Ivănescu Asoc.Prof. Dr. Eng. Sanda Levcovici AFFILIATED WITH: - ROMANIAN SOCIETY FOR METALLURGY - ROMANIAN SOCIETY FOR CHEMISTRY - ROMANIAN SOCIETY FOR BIOMATERIALS - ROMANIAN TECHNICAL FOUNDRY SOCIETY - THE MATERIALS INFORMATION SOCIETY (ASM INTERNATIONAL)


FASCICLE IX METALLURGY AND MATERIALS SCIENCE, ISSN 1453 – 083X NR 1 – 2004

Table of content

1. Constantin Baciu, Maria Baciu, Mihai Lozovan – About wear behaviour of super-alloyed steels in electrolytic plasma…………………………………………… 2. N. Grechanyuk, V. Osokin, P. Shpak, I. Grechanyuk – Equipments and application of electron-beam technology for obtaining of new types of materials, coatings and re-melting………………………………………………………………. 3. Dan Teodor Levcovici, Sanda Maria Levcovici, Adriana Preda – Alloying and dispersion of hard particles into titanium and titanium alloy………………………… 4.. Vasile Marina – Reading the possibilities to decode the microstructure characteristics from macroexperience ……………………………………………….. 5. Nicolae Cănănău, Alexandru Ivănescu, Lilica Ivănescu – Simulation of semi-planetary rolling process at experimental stand……………………………………… 6. Tamara Radu, Olga Mitoşeriu, Lucica Balint – Experimental researches on the Zn–Fe coating………………………………………………………………………… 7. Adriana Preda, Dionisie Bojin, Sanda Maria Levcovici – Researches concerning heat treatment effects upon the morphology of high-tensile foundry brass structure………………….…………………………………………………………… 8. Dumitru Dima - Obtaining and characterization of ferrite powders from metallurgycal waste materials………………………………………………………… 9. Petrică Alexandru – The plastic deformation of steel sheets for enameling and defects of enameled layer……………………………………………………………. 10. Simona Boiciuc – Surface hardening by means of microalloying and vibratory electrode deposit……………………………………………………………………… 11. Ionel Petrea – Mathematical model for the sintering iron-graphite powders mixture simulation……………………………………………………………………. 12. Doru C. Hanganu – Cross-section modification, pending open die forging process…………….………………………………………………………………….. 13. Maria Vlad, Olga Mitoşeriu, Emil Strătulat – Nonferrous alloys with special properties in high temperature………………………………………………………... 14. Ciurea Aurel, Bordei Marian, Eftimie Dorin – Mathematical modeling of rotation în case of vertical centrifual casting…………………………………………. 15. Petrică Alexandru - The hardening of thin bands by plain carbon steel with low content of carbon by quenching….…………………………………………………… 16. Stela Constantinescu, Olga Mitoşeriu - Studies and researches to improve the structure and hardness of steel C420………………………………………………….

5 8

11

19

26

31

36

43

49

54

62

67

70

73

77

82



5

ABOUT WEAR BEHAVIOUR OF SUPER-ALLOYED

STEELS IN ELECTROLYTIC PLASMA

Constantin BACIU1, Maria BACIU1, Mihai LOZOVAN2,

1Gh.Asachi” Technical University of Iaşi, 2National Institute for Research-Development in Technical Physics, Iaşi

email: [email protected]

ABSTRACT

Considering OLC55 and 40Cr10 steels, nitrided and quenched in electrolytic plasma, their wear resistance was expressed by means of the intensity (Im) and rate (Vm) of mass wear. The thermal processing variants applied to the two steels under study comprised: three different diffusion temperatures, Td = 650, 700 and 7500C; two values of the diffusion time, td = 3 and 6 min; the electrolyte with the composition: 10 % NH4Cl + 20 % NH4OH + H2O. Based on the experimental data recorded for the two steels under study, the variation curves were plotted, as a function of the diffusion temperature, for the

average mass wear rate, mV = f(Td) and for the average mass wear intensity, mI = f(Td).

KEYWORDS: electrolytic plasma, steels, abrasive wear

1. Introduction

Microstructural changes caused by the process of super-alloying by diffusion during steel heating in electrolytic plasma influence wear resistance. Mass wear of materials can be expressed by means of the mass wear rate, Vm:

Vm = t

cmΔ, [g/h] (1)

and by the mass wear intensity, Im:

Im = ,fLfA

m

⋅

Δ [g/m3] (2)

where; Δmc represents cumulated mass loss, Δm - mass variation, t – time, Af - friction surface and Lf – length of the friction path.

2. Experimental

Wear behaviour of OLC55 and 40Cr10 steels, super-alloyed with nitrogen by heating in electrolyte solution has been studied under the conditions of a dry wear regime between the contact surfaces of the specimens and the diamond disk, Figure 1.

Fig. 1. Conditions specific to the abrasive wear test (Vslip = 1.25 m/min, p = 1.5 MPa).

The technological variants of the nitriding treatments applied to OLC55 and 40Cr10 steels are shown in Figure 2. The specimens, subjected to the wear tests, were super-alloyed with nitrogen in electrolytic plasma according to the conditions given in Table 1.

mailto:[email protected]



6

Fig. 2 Technological parameters of the nitriding treatments by heating in electrolytic plasma:

variant a – nitriding (Td = 6500C, td = 3; 6 min) and quenching; variant b – nitriding (Td = 6500C, td = 3; 6 min) followed by heating for austenitization (Taust = 7500C, t = 10 sec)

and quenching; variant c – nitriding (Td = 7500C, td = 3; 6 min) followed by a quenching step (6000C, t = 30 sec) then heating for austenitization (Taust = 7500C, t = 10 sec) and quenching;

Table 1

Technological parameters of

diffusion No. Steel grade Marking

Td, [0C] td, [min]

Initial mass, mi, [g]

1 3R 650 40.3547 2 3I 700 40.4266 3

40Cr10 3M 750 40.5210

4 4D 650 40.7285 5 4F 700 40.7407 6

OLC55 4S 750

3

41.1910 7 3A 650 40.6015 8 3V 700 39.1349 9

40Cr10 3J 750 40.6167

10 4Y 650 40.3947 11 4DD 700 40.9206 12

OLC55 4T 750

6

40.4801 Observations:

- electrolyte: 10 % NH4Cl + 20 % NH4OH + H2O - after nitriding all specimens were quenched.

3. Results

Measurements of the mass losses, Δm and Δmc, were performed by weighing the specimens on the analytical balance, at different time intervals (1; 2; 2.5; 3; 3.5; 4, 4.5 and 5 hours) of the test. Based on the relationships (1) and (2) the values of mass wear rate and mass wear intensity, respectively, were determined and subsequently the

average values of these parameters were calculated, (Table 2).

The variation of the two mass wear parameters, mV and mI , as a function of the diffusion temperature, is shown in Figure 3 and 4, respectively, for the two steels under study.

Table 2

Specimen Average values 3R 3I 3M 3A 3V 3J 4D 4F 4S 4Y 4DD 4T

mV 8.66 12.71 18.06 8.11 14.38 19.23 18.66 18.98 20.82 11.56 13.72 20.11

mI 6.503 10.802 13.568 6.094 9.544 14.445 14.014 14.252 15.633 8.683 10.301 15.105



7

Fig. 3. Variation of the average mass wear rate, v m = f(Td), as a function of diffusion temperature:

Fig.4. Variation of the average mass wear intensity, I m = f(Td), as a function of diffusion temperature:

⎩⎨⎧

−

−=

55OLC3

10Cr4013mint

⎩⎨⎧−−

=55OLC310Cr401

3mint

⎩⎨⎧

−

−=

55OLC4

10Cr402min6t

⎩⎨⎧

−−

=55OLC410Cr402

min6t

4. Conclusion

i) Nitrogen super-alloying of steels, heated in

electrolytic plasma provides considerable improvement of wear behaviour, in dry friction regime.

(ii) Under the condition of maintaining constant the diffusion time, the lowest values of the parameters of wear process correspond to Td = 6500C; the increase of diffusion temperature causes the gradual increase of mass losses Δm and Δmc.

(iii) By comparing the above two values of the diffusion time, better wear behaviour has been noticed for the specimens which were thermally processed for 6 minutes.

(iv) The analysis of the experimental data emphasizes the superior wear behaviour of 40Cr10

steel as compared to OLC55 steel under the conditions of the same values for the parameters of diffusion process.

References

[1]. Aleksandrov, V.N. Ignat'kov, D.A., Pasinkovskÿ, E.A., Fizico-mehaničeskie svojstva stali 45, azotirovannoj v èlektrolitnoj plasme, Elektronnaâ obrabotka materialov, nr. 2, p.17-18, 1982. [2]. Baciu Maria Contributions on the structural and property changes of thermally and thermochemically treated steels in electrolytic plasma (in Romanian), PhD thesis, „Gh.Asachi” Technical University from Iaşi, 1999. [3]. Belkin, P.N, Ignat'kov, D.A., Pasinkovskÿ, E.A., Azotirovanie v elektrolitnoj plazma, Kolloquium Eigensspannungen und Oberflächen-verfestigung, p.265, 1982.



8

EQUIPMENTS AND APPLICATION OF ELECTRON-BEAM

TECHNOLOGY FOR OBTAINING OF NEW TYPES OF MATERIALS, COATINGS AND RE-MELTING

N. GRECHANYUK, V. OSOKIN, P. SHPAK, I. GRECHANYUK

Scientific-and-Producing Enterprise "GEKONT"

Ukraine, Vinnitsa, Vatutina str., build. 25

ABSTRACT

In this reviews modern ways, examples of operational use of electron-beam technologies are shortly submitted. Also are described of its advantage and prospect for its development. In the article examples of successful operational use of EB-technologies in applied materials science and industry are adduced.

KEYWORDS: electron-beam technology, composite material, electrocontact

material, protective coating, electron-beam re-melting.

1. Introduction

The development of economics of any state bodily depends on a level of development of science and application of high and modern technologies. One of newest, dynamically developing and perspective technologies widely used in applied materials technology, is the electron-beam technology.

Modern concept the electron-beam technology are includes:

- electron-beam melting; - electron-beam evaporation (EB-PVD) and

subsequent condensation in vacuum of metallic and non-metallic materials by covers or massive bars, unbound from a substrate, of composite materials, powder and etc.;

- electron-beam welding; - electron-beam treatment. The results of experimental researches and

industrial approbation some of these technologies materials and coatings are below submitted.

2.Materials for electrical contacts

One of the most perspective applications of an electron-beam method can be the creation on its base of new technologies of formation of materials for producing of electrical contacts. Now in all advanced industrial countries of the world as materials for electrical contacts, as a rule, are used alloys on the basis of noble metals, first of all, alloys on the basis of silver with the additives of cadmium oxides, tin, tungsten, nickel, which are received by the methods of powder metallurgy.

Taking into account constantly growing demands and prices for silver, tungsten and other metals, the researchers all over the world purposefully works for creation cheaper, without a decrease of the operational characteristics of pseudo-alloys on the basis of silver, composites for electrical contacts not containing noble metals.

In the "GEKONT"-company (Ukraine) the idea of production of micro-layered materials was realized, created and patented [1] by means of the industrial electron-beam technology of production of materials for electrical contacts, which have no having of analogues in world practice.

Fig. 1. The scheme of industrial EB-installation for production of micro-layer materials.





9

Composite condensed materials Cu - (0,08…0,2)mas.% Zr - (0,08…0,2) mas.% Y -(6…14) mas.% Mo was received on industrial electron-beam installation UE-189 as a sheet by a diameter of 800 mm and thickness from 0,5 up to 4 mm. The scheme of the installation is shown in a Fig. 1.

The typical structure of micro-layer materials used for electrical contacts consists from alternating layers: low-doped alloy on a basis of copper with thickness 0,5-0,8 μm and molybdenum layers with thickness 0,15-0,2 μm. The stable layered structure is kept up to 900°C. Density of a material – 9,07 g/sm3; strength σs in a longitudinal direction is 600-900 MPa; strength σs' in a cross direction is 500-700 MPa, relative elongation is 2-9 %; electrical conductivity – 0,027 Ωm⋅mm2/m.

The materials are well treated by cutting, extrusion, drilling, are easily soldered by any of known ways of soldering with usage standard argentums and non-argentums of solders.

To the present time it is made more than 1.5 million contacts of 302 types.

In a stage of industrial approbation there are also contacts of system copper-chromium for arc-extinguish chambers [2].

3.Porous materials and coatings for

filters and catalysts

The porous materials are widely used for clearing of impurity of liquids and gases. The areas of their application constantly extend: in space engineering they are used as elements of ion engines, and also for cooling by evaporating; in aircraft - for prevention of plane icing, in transport - for production of an electrical current etc.

Until recently products from porous materials were produced only by methods of powder metallurgy.

In "GEKONT"-company the new way of formation of porous materials is developed [3], which includes evaporation from separate water-cooled crucibles of metals or alloys and non-metal compounds and simultaneous their deposition on the substrate, previously heated up and covered with a dividing l ayer, on which there is a branch of a porous material from a substrate. The regulation of the average size in a range from 0,1 μm up to 1 μm and volume in a range from 5 up to 40 % is carried out by change of concentration of non-metal compounds, and also regulation of speed and temperature of heating of the deposited material. By the given technology it is easy to receive porous materials on the basis of pure metals and metal alloys with high mechanical characteristics. For example, the condensate of pure titanium without porous, received at temperature of deposition 620°C has the following

mechanical characteristics: σs=540 MPa; σ02=430 MPa; δ=22 %. A porous condensate of pure titanium with volume of porous 19.5 % (volume of open porous 12,7 %) with the average size of porous 0,4-0,5 μm, annealed by a special mode in vacuum have the following properties: σs=830 MPa; σ02 =720 MPa; δ=4.5 %.

4. Protective coatings for the gas turbines blades

In "GEKONT"- company the coatings on the

blades of gas turbines produced by electron- beam evaporation of alloys MeCrAlY (where Me- Ni, Co, Fe), MeCrAlYHfSiZr and ceramics on a basis ZrO2, stabilized Y2O3 and subsequent condensation of a vapour phase on a working and directing blades of gas turbines of various purpose [4, 5].

Three types of coatings are developed: -one-layer metal-alloy such as MeCrAlY,

MeCrAlYHfSiZr, composite of micro-layer such as with alternation of layers MeCrAlY, MeCrAlYHfSiZr and MeCrAlY + МеО, MeCrAlYHfSiZr + MeO, where MeO - Al2O3 or ZrO2 + 6-8%Y2O3;

-two-layer coatings with internal metal MeCrAlY, MeCrAlYHfSiZr and external ceramic ZrO2-Y2O3 layers; two-layer coatings with internal composite MeCrAlY + МеО, MeCrAlYHfSiZr + MeO and external ceramic ZrO2-Y2O3 layers;

-three-layer of a coating with internal metal-alloy MeCrAlY, MeCrAlYHfSiZr, intermediate composite MeCrAlY + MeO, MeCrAlYHfSiZr + МеО dispersion-strengthened or micro-layer types and external ceramic ZrO2-Y2O3 layer.

The EB-PVD of protective coatings to the TBC - implements on one type of equipment and during one technological cycle.

5. Electron-beam re-melting of

metals and alloys

Electron-beam re-melting is widely applied in the world practice to producing metals high-melting, high-reactionary and precision alloys of high quality and purity. It is one of the most effective ways of increase of metal purity from parasitic and non-metallic impurity and also elimination of chemical and structural non-uniformity.

The specified technology has the following advantages:

- re-melting process and formation of ingots occur for one working cycle without subsequent heat- and –mechanical treatments of (hammering, annealing);

- an opportunity re-melting of chunks wastes;



10

- fast change of equipment for manufacturing ingots with necessary sizes;

- high quality of ingots after vacuum re-melting; - production of small parties of commodity ingots. Electron-beam re-melting also produce ingots Mo,

Nb, Hf, Zr, Ni, Cu, Co, Fe, and super-alloys on their basis, intermetallics Ti3Al, TiAl, Ni3Al, NiAl.

In "GEKONT"-company the industrial technology of electron-beam re-melting of wastes of high-speed steels was mastered and production of commodity ingots for the subsequent manufacturing of the high quality of cutting tools from them [6, 7] was developed.

Sizes of preparations: - ingots - diameter from 60 up to 138 mm, length

up to 1900 mm; - slabs - section 140 × 160 mm, length up to 1900 mm.

6. The laboratory and industrial electron-

beam equipment for realization of the above mentioned technological processes

By “GEKONT”-company have designed, made,

delivered and installed to the customers follow EB-installations [8]:

- to the "Zarya"- turbine enterprise, Nikolaev-city, Ukraine. Industrial electron- beam facility L-1 for depositing of coatings on the gas turbines blades. The sizes of covering blades: a diameter up to 350 mm, length up to 700 mm;

- to the Harbin Institute of Technology (HIT), Harbin, China. Industrial electron-beam facility for prodction of composite materials with productivity up to 10 tons per one year;

- to the Korean Institute of Science and Technology (KIST) of Seoul, Korea. Laboratory electron-beam facility L-2, allowing to receive composite materials as a disk by a diameter of 500 mm and thickness up to 5 mm;

- to deposits coatings on turbines blades. The sizes of coated blades: - a diameter up to 220 mm, length up to 500 mm.

Production of ingots of metals and alloys by a

diameter from 70 mm up to 100 mm and length up to 500 mm.

At the enterprise also the electron-beam guns and industrial electron-beam equipment of new generation were designed for deposition of functional graded coatings and production of ingots of metals and alloys with productivity up to 300 tons per one year.

References

[1]. N. Grechanyuk, V. Osokin, I. Afanasyev, I. Grechanyuk The composite material for electrical contacts and method of their production./ the Patent of Ukraine #34875 from 16.12.2002. [2]. N. Grechanuyk, N. Plaschenko, V. Osokin The contact material for arc-extinguish of chambers and method of their reception./ the Patent of Ukraine #32368A from 15.12.2000. [3]. N. Grechanyuk, V. Osokin, I. Afanasyev, I. Grechanyuk , Y. Piyuk The method of production of porous materials./ the Patent of Ukraine #46855 from 17.06.2002. [4]. N. Grechanyuk, P. Kucherenko, V. Osokin, P. Shpak The modern condition and prospects of creation of heat-protective coatings (TBC) for gas-engines blades and equipment for their production./ The news of energetic, #9, 2000, pp. 32-37, Kiev. [5]. N. Grechanyuk, P. Kucherenko, V. Osokin, I. Afanasyev, S. Belik, V. Akrimov, I. Grechanyuk, Y. Piyuk The coatings for the gas-engine blades./ the Patent of Ukraine #42052 from 15.10.2001. [6]. N. Grechanyuk, P. Kucherenko, V. Osokin The method of manufacturing of rotating bodies by a method level-by-level solidification and device for their production./ the Patent of Ukraine # 18135A from 31.10.1997. [7].N. Grechanyuk, I. Afanasyev, P. Shpak, V. Osokin, M. Shwedchikov The method of manufacturing of bars for the tool from high-speed steel and device for their production./ the Patent of Ukraine #37658A from 15.05.2001. [8].N. Grechanyuk, P. Kucherenko Facility for electron-beam depositing of coating./ the Decision on distribution by the patent of Ukraine on the invention under the application #99116199 from 15.11.1999.



11

ALLOYING AND DISPERSION OF HARD PARTICLES

INTO TITANIUM AND TITANIUM ALLOY

Dan Teodor LEVCOVICI1, Sanda Maria LEVCOVICI2, Adriana PREDA1

1 S.C.”UZINSIDER ENGINEERING”-S.A.,2 Smardan, 800701 Galatzi, Romania, 2 ”DUNĂREA DE JOS” University of Galatzi, 47 Domneasca, 800008 Galatzi, Romania,

[email protected].

ABSTRACT

Titanium and its alloys exhibits several very good properties, e.g. a high strength-to-weight ratio and an excellent corrosion resistance. However, the tribological properties of titanium are relatively poor. In order to improve the wear resistance of Ti-6Al-4V alloy (Grade 5, ASTM B 265-94) and unalloyed titanium (Grade 2, ASTM B 265-94), the specimens of 35x35x15 mm3 were two-step surface alloyed using added materials with hard particles of carbides (WC, SiC), borides (TiB2) and graphite for comparison. The simultaneous melting of added material and surface layer by a CO2 continuous wave laser. For the laboratory experiments the light optical microscope with automatic image analysis, X-ray diffractometry, microhardness meter, etc., were used. KEYWORDS: Laser, Local Melting, Surface Alloying, Layer, Hard Particles.

1. Introduction

In many industrial plants for the processing and

storing of the fluid materials the hydrodynamic transport processes were increased for economical reasons. The question arises that the parts made of titanium and its alloys should present not only suitable corrosion strength, but also a higher strength to the abrasive action of the complex environment where they operate. It is known that solid-state heat treatment processes cannot harden the titanium. Therefore research works were conducted to determine the most effective methods of surface hardening by local melting, [1] - [6]. In this report the results of the research works performed for the surface hardening of the Grade 5 titanium alloy and Grade 2 titanium by simultaneous melting of the surface layer and

several added materials with hard particles of carbides (WC, SiC), borides (TiB2) and graphite for comparison are shown. The local melting was performed using a CO2 continuous wave laser provided with a numerically controlled x-y-z table.

2. Material and Sample Preparation

The experiments for the determination of the hardening conditions of titanium alloy Grade 5 (BM 1 - base material, with 264 HV49) and of titanium Grade 2 (BM 2 - base material, with 262 HV49) were made on plates of 35x35x15 mm3. The added material (AM) was made as a paste by mixing a part of hydroxiethyl cellulose with different parts of powder. The composition and specific weight of the added material are shown in Table 1

.

Table 1. Added material composition. BM

code Added

material AM weight

[mg/cm2] Composition in % weight

5 13.2 100 - - - - - 6 84.0 - 100 - - - - 7 46.5 - - 85 5 10

Ti

(BM 2) 8 34.0 - 33 - - - 67






12

Table 2. Laser working regime codes.

Working regime code Working parameters 1 2 3 4 5 6

Power P [W] 1300 Laser beam diameter on the specimen surface [mm] 1.8 (1.65 for specimen code 4)

Laser-scanning rate [mm/s] 25 22 19 16 13 11

Table 3. The metallografic analysis and the hardness measurements for Ti-6Al-4V alloyed by laser. Laser

parameters Surface layer

geometry [mm] Alloyed layer Hardened layer

Test code v

[mm/s] K

[J/mm2] WA DA WH DH

HV49 [MPa]

Microstructural characteristics

1.1 25 28.9 2.25 0.39 3.05 0.90 469 M+E+HP 1.2 22 32.8 2.45 0.60 3.25 1.15 450 M+E+HP 1.3 19 38.0 2.68 0.64 3.35 1.24 466 M+E+HP 1.4 16 45.1 2.75 0.62 3.55 1.35 466 M+E+HP 2.2 22 32.8 2.64 0.50 3.22 1.04 466 M+E+HP 2.3 19 38.0 2.40 0.49 3.14 1.00 501 M+E+HP 2.4 16 45.1 2.40 0.48 3.10 1.00 490 M+E+HP 2.5 13 55.5 2.50 0.49 3.25 1.10 549 M+E+HP+P 2.6 10 65.7 2.35 0.42 3.20 0.99 495 M+E+HP+P 3.2 22 32.8 - - - - 462 M+E+HP 3.3 19 38.0 2.60 0.43 3.50 1.22 418 M+E+HP 3.4 16 45.1 2.60 0.65 3.45 1.10 415 M+E+HP 3.5 13 55.5 2.44 0.45 3.38 1.07 456 M+E+HP 3.6 10 65.7 - - - - 450 M+E+HP+P 4.2 22 35.8 2.20 0.30 2.88 0.90 502 M+E+HP 4.3 19 41.5 2.30 0.28 2.95 0.95 498 M+E+HP 4.4 16 49.2 2.38 0.45 3.0 0.99 564 M+E+HP 4.5 13 60.6 2.40 0.47 3.25 1.02 549 M+E+HP

Table 4. The metallografic analysis and the hardness measurements for Ti alloyed by laser.

Laser parameters

Surface layer geometry [mm]

Alloyed layer Hardened layer

Test code v

[mm/s] K

[J/mm2] WA DA WH DH

HV49 [MPa]

Microstructural characteristics

5.1 25 28.9 - - - - 381 M+E+HP 5.2 22 32.8 2.20 0.48 3.40 1.20 311 M+E+HP 5.3 19 38.0 2.25 0.55 3.45 1.28 347 M+E+HP 5.4 16 45.1 2.15 0.26 - 1.80 330 M+E+HP 6.2 22 32.8 1.95 0.28 2.70 0.82 458 M+E+HP 6.3 19 38.0 1.84 0.36 2.60 0.82 452 M+E+HP 6.4 16 45.1 2.00 0.36 2.83 0.83 458 M+E+HP 6.5 13 55.5 2.20 0.46 2.88 0.92 463 M+E+HP+P 6.6 10 65.7 2.15 0.43 2.80 0.82 463 M+E+HP+P 7.2 22 32.8 2.30 0.30 3.20 0.88 487 M+E+HP 7.3 19 38.0 2.45 0.40 3.20 1.00 490 M+E+HP 7.4 16 45.1 2.46 0.47 3.40 1.00 516 M+E+HP 7.5 13 55.5 2.56 0.54 3.48 1.19 487 M+E+HP 7.6 10 65.7 2.60 0.55 3.60 1.20 480 M+E+HP+P 8.1 25 28.9 2.18 0.34 - 0.88 288 M+E+HP 8.2 22 32.8 2.30 0.30 3.15 0.92 399 M+E+HP 8.3 19 38.0 2.40 0.34 3.22 0.93 423 M+E+HP 8.4 16 45.1 2.54 0.37 3.40 0.98 444 M+E+HP 8.5 13 55.5 2.70 0.47 3.55 1.03 444 M+E+HP



13

For experiments a CO2 continuous wave laser was

used. The radiation power P=1300 W, the laser beam diameter on the specimen surface ds=1.8 mm (by a lens of 127 mm focal distance) and laser-scanning rate v=10-25 mm/s were used. The working regime by energy factor K=P(v·ds)-1 was characterised. After paste covering and drying, the specimens were fitted on a digitally controlled x-y-z table, coupled with a CO2 continuous wave laser. The paste layer and the base material were melted simultaneously, under argon protection. The laser working regime codes are shown in Table 2. The surface layers thus obtained were characterized by X-ray diffractometry and HV49 (49N load) hardness measurements. The micro-structure characteristics, width, depth and micro-hardness HV0.98 (0.98N load) of the alloyed zone on the specimen cross section, metallographically prepared, by optical microscopy were determined.

3. Results and discussions

The working parameter values, the surface

layer geometry, the metallographic analysis and the hardness measurements made on single bands processed with different laser scanning rates are shown in Table 3 and 4, (WA, DA - alloying layer width and depth, respectively, WH, DH - hardening layer width and depth, respectively, M - martensite,

E - interdendritic eutectic, HP - hard particles, P - pores). General views of the microstructure in the cross section of the laser alloyed layers for 1 - 8 added materials, respectively are shown in Figures 1-6. The microstructure of the laser alloyed layer present a dendritic columnar aspect. The ultrarapid thermal cycle in the Ac1 – solidus temperature range provides the martensitic hardening. In specimens 1.2 and 5.2, under conditions of carbon alloying, it consist of ultrafine martensite, oversaturated, and secondary carbides. In specimens 2.2, 3.2, 4.2, 6.2, 7.2, 8.2, under conditions of complex alloying, ultrafine martensite, oversaturated by partial carbide dissolving, chemically non-homogenous and a developed substructure is formed. In Figures 7 a-d, the variation of depth of the laser-alloyed layers with the density energy factor, on the BMs 1 and 2, respectively, are shown. While the 1 added material on the BM 1 provides the maximum alloying depth above 33 J/mm2, and 3 added material on the BM 1 for 45 J/mm2, 4 added material on the BM 1, the 6, 7, 8 added materials on the BM 2 up to this density energy factor, provides the minimum alloying depth. The variation of the relative maximum intensity of the phases of the laser-alloyed layers with the diffraction angle for specimens 2.2, 6.2, 3.2, 7.2, 4.2, 8.1 in Figures 8 - 13, are shown.

Fig. 1. The microstructure of the laser hardened titanium.

Fig. 2. The microstructure of the laser alloyed layer for 1.2

specimen.


specimen.


specimen.

Fig. 5. The microstructure of the laser alloyed layer for 4.2 specimen


specimen.



14

00.20.40.60.8

20 30 40 50density energy factor K [J/mm2]

allo

yed

laye

r dep

th

[mm

]

1 5

00.20.40.60.8


allo

yed

laye

r dep

th

[mm

]

2 6

a. b.

00.20.40.60.8


allo

yed

laye

r dep

th

[mm

]

3 7

00.20.40.60.8


allo

yed

laye

r dep

th [m

m]

4 8

c. d.

Fig. 7 a-d. Variation of depth of the laser-alloyed layers with the density energy factor, on the BMs 1 and 2.

Ti-2-rtf

Unghi de difractie 2 [ ]

I/Imax

., [%

]

0

10

20

30

40

50

60

70

80

90

100

30 40 50 60 70 80 90 100 110 120 130 140

code 2.2

Ti (1

01)

TiC (2

00)

SiC

(004

)TiC

(111

)Ti

(002

)

Ti (1

02)

SiC

(103

) TiC (2

20)

Ti (1

10)

SiC

(243

)

TiC (4

00)

SiC

(110

)

Ti (2

01)

SiC

(201

)

SiC

(210

)TiC

(331

)

Ti (1

03)

θ o

Diffraction angle [o]

Fig. 8. Variation of the relative maximum intensity of the components with the diffraction angle for specimen 2.2.

For diffractometric analysis of the alloyed layers obtained for each added and base material, respectively, the samples by multistep scanning with laser beam at a scanning rate v = 22 mm/sec and a 35% overlapping band structure was per-formed.

The working parameters: U=30kV, I=32mA, ω=4o/min, KαCo(λ=1.7902066Å), 2; 2; 0.5, 2θI=40o, Δθ= 90o.



15

Ti-3-rtf code 3.2


I/Imax

., [%

]

0

10

20

30

40

50

60

70

80

90

100

30 40 50 60 70 80 90 100 110 120 130 140

WC

(001

)

TiC (1

11)Ti

(002

)

Ti (1

01)

Ti (1

02)

SiC

(004

)

WC

(101

)TiC

(200

)Si

C (1

03)

WC

(110

) SiC

(201

)W

C (1

11)

WC

(100

)

WC

(201

)

Ti (2

01)

TiC (4

00)

WC

(112

)Si

C (2

10)

Ti (1

03)

TiC (3

31)

oθ

Fig. 9. Variation of the relative maximum intensity of the components


with the diffraction angle for specimen 3.2.

Ti-4-rtf


I/Imax

., [%

]

0

10

20

30

40

50

60

70

80

90

100

30 40 50 60 70 80 90 100 110 120 130 140

θ o

SiC

(004

)TiB

2 (10

0)TiC

(111

)Ti

(002

)

Ti (1

01)

TiC (2

00)

TiB2 (

101)

Ti (1

02)

SiC

(110

)TiC

(220

)Ti

(110

) TiB2 (

102)

Ti (1

03)

TiC (3

11)

SiC

(203

) Ti (2

01)

TiB2 (

112)

TiC (4

00)

TiB2 (

202)

SiC

(210

)TiC

(331

)

code 4.2


Fig. 10. Variation of the relative maximum intensity of the components with the diffraction angle for specimen 4.2.

Ti-7-rtf


I/Imax

., [%

]

0

10

20

30

40

50

60

70

80

90

100

30 40 50 60 70 80 90 100 110 120 130 140

θ o

Ti (3

31)

SiC

(210

)W

C (1

12)

SiC

(243

)

Ti (2

01)

WC

(111

)Ti

(103

)

WC

(002

)Ti

(110

)

Ti (1

02)TiC

(200

)Ti

(101

)

Ti (0

02)

WC

(100

)Si

C (0

04)

TiC (1

11)

code 7.2 code 6.2






16

Ti-7-rtf


I/Imax

., [%

]

0

10

20

30

40

50

60

70

80

90

100

30 40 50 60 70 80 90 100 110 120 130 140

θ o

Ti (3

31)

SiC

(210

)W

C (1

12)

SiC

(243

)

Ti (2

01)

WC

(111

)Ti

(103

)

WC

(002

)Ti

(110

)

Ti (1

02)TiC

(200

)Ti

(101

)

Ti (0

02)

WC

(100

)Si

C (0

04)

TiC (1

11)


code 7.2



Ti-8-rtf


I/Imax

., [%

]

0

10

20

30

40

50

60

70

80

90

100

30 40 50 60 70 80 90 100 110 120 130 140

θ o

code 8.1

Ti (1

01)

SiC

(004

)TiC

(111

)Ti

(002

)TiC

(200

)TiB

2 (10

1)Ti

2B (3

10)

Ti (1

02)

SiC

(110

)TiC

(220

)

Ti (1

12)

TiC (2

22)

Ti2B

(332

)

TiB2 (

102)

TiB2 (

112)

TiB2 (

202)

SiC

(210

)

TiC (4

00)

Ti (1

10)




In Figures 14 a, b and 15 a, b, the variation of HV49 hardness of the laser-alloyed layers with the density energy factor, on the BMs 1 and 2, respectively, are shown. In Figures 15 and 16 the microhardness variation with the laser-alloyed layers depth for each added material and base material are shown.

While the 1, 2, and 4 added materials on the of Ti-6Al-4V alloy provides the maximum hardness, the 6 added material on the unalloyed titanium provides the maximum hardness, for K = 38 – 45 J/mm2. The most important increments of the HV49 hardness were achieved by laser alloying of Ti-6Al-4V alloy, using the 4 added material, for K = 45 – 61 J/mm2. Also,

high values of hardness were obtained as a result of using the2 added material at the same base material, K = 55 J/mm2.

The 1, 6, and 7 added materials caused the lowermost variation of the hardness on the alloyed surface of the 1 and 2 base material, respectively, with the density energy factor. The 2, 4 and 8 added materials provided the highest variation of the hardness on the alloyed surface of the 1 and 2 base material, respectively, with the decrease of the laser-scanning rate.

In any case the microhardness variation (Figures 16 and 17) on the depth of the processed layer points out the higher hardening of the alloyed layers. The



17

maximum microhardness values are found in the special 3, 6 added materials, for various density energy values.

The analysis on the light optical microscope showed that all added materials led to compact layers.

However, at low scanning rates the layers alloyed with the 2, 3, 6, and 7 added materials showed pores. Such defects were not present in all the layers alloyed with the 1, 4, 5 and 8 added materials.

200300

400500

600

20 30 40 50

density energy factor K [J/mm2]

hard

ness

HV

49

[MPa

]

1 5

a.

200

300

400

500

600

20 30 40 50 60 70


hard

ness

HV

49

[MPa

]

2 6

b. Fig. 14 a, b. Variation of HV49 hardness of the laser-alloyed layers with the density energy factor,for AM 1, 2

on the BM 1 (a) and AM 5, 6 on the BM 2 (b) respectively.

200

300

400

500

600

20 30 40 50 60 70


hard

ness

HV

49

[MPa

]

3 7

a.

200300400500600

20 30 40 50 60 70


hard

ness

HV

0.4

9 [M

Pa]

4 8

b. Fig. 15 a, b. Variation of HV49 hardness of the laser-alloyed layers with the density energy factor, for AM 3, 4

on the BM 1and AM 7, 8 on the BM 2 respectively.

010002000300040005000600070008000

0 0.5 1 1.5depth [mm]

mic

roha

rdne

ss H

V 0

.98

[MPa

]

code 1.1code 1.2code 1.3code 1.4

0

1000

2000

3000

4000

5000

6000

7000

8000

0 0.5 1 1.5depth [mm]

mic

roha

rdne

ss H

V 0

.98

[MPa

]

code 2.2code 2.3code 2.4code 2.5code 2.6

010002000300040005000600070008000

0 0.5 1 1.5depth [mm]

mic

roha

rdne

ss H

V 0

.98

[MPa

]

code 3.2code 3.3code 3.4code 3.5

a. b. c.



18

010002000300040005000600070008000

0 0.5 1 1.5depth [mm]

mic

roha

rdne

ss H

V 0

.98

[MPa

]code 4.2code 4.3code 4.4code 4.5

MB 1

0

1000

2000

3000

4000

5000

6000

7000

8000

0 0.5 1 1.5

depth [mm]m

icro

hard

ness

HV

0.9

8 [M

Pa] code 5.1

code 5.2code 5.3code 5.4

MB 2

0100020003000400050006000700080009000

0 0.5 1 1.5

depth [mm]

mic

roha

rdne

ss H

V 0

.98

[MPa

] code 6.2code 6.3code 6.5code 6.6

MB 2 d. e. f.

Fig. 16 a - f. Microhardness variation with the laser-alloyed layers depth for AM 1, 2, 3 and 4 on the BM 1and AM 5, 6 on the BM 2 respectively.

010002000300040005000600070008000

0 0.5 1 1.5depth [mm]

mic

roha

rdne

ss H

V0.

98 [M

Pa] code 7.1

code 7.3code 7.4code 7.5

0

1000

2000

3000

4000

5000

6000

7000

8000

0 0.5 1 1.5

depth [mm]

mic

roha

rdne

ss H

V 0

.98

[MPa

] code 8.1code 8.2code 8.3code 8.4code 8.5

a. b.

Fig. 17 a, b. Microhardness variation with the laser-alloyed layers depth for AM 7 and 8 on the BM 2.

4. Conclusion

By laser surface alloying using the two-steps

procedure with the different added materials consisting of hydroxiethyl cellulose and powder with hard particles of carbides (WC, SiC), borides (TiB2) and graphite for comparison, continuous and smooth alloyed layers can be obtained, showing an average hardness 1.5 - 2 higher than the titanium and its alloy (Grade 2, Grade 5, ASTM B 265-94).

References [1]. Weisheit, A., Mordike, B.L., 1992, “Lasers Surface Alloying of Titanium with Nickel and Cobalt”, Proc.Eur. Conf. on LaserTreatment of Materials ECLAT’92, by B.L. Mordike, Göttingen, pp. 229. [2]. Villechaise, P., Peraud, S., Mendez, J., 1995, “Fatigue Resistance Improvement of a 316L Stainless Steel and Ti-6Al-4V

Titanium Alloy by Dynamically Ion Mixed NiTi Coatings”, Proc. 8th Int. Conf. Surface Modification Technologies SMT VIII, by T.S. Sudarshan and M. Jeandin, London, pp.461. [3]. Levcovici, S.M., Levcovici, D.T., Gologan, V., FARKAŞ, L., 1997, “Contributions to the characterization of the WC – steel composite structures, obtained by laser melting”, Proc. 5th Eur. Conf. Advanced Mat. and Proc. EUROMAT ′97, by L.A.J.L. Sarton and H.B. Zeedijk, Zwijndrecht, pp.131. [4]. C. Langlade, A.B. Vannes, 1998, Characterisation and Applications of Laser Treated Titanium and Titanium Alloys”, Proc. 11th Int. Conf. Surface Modification Technologies SMT XI, by T.S. Sudarshan, M. Jeandin and K.A. Khor, London, pp.634. [5]. Levcovici, D.T., Munteanu, V., Levcovici, S.M. Mitoseriu, O., Benea, L., Paraschiv, M.M., 1999, “Laser Processing of MMC Layers on a Metal Base”, Materials and Manufacturing Processes, Vol. 14, No. 4, by T.S. Sudarshan and T.S. Srivatsan, New York, pp.475. [6]. S.M. Levcovici, D.T. Levcovici, V. Munteanu, M.M. Paraschiv, A. Preda, 2000, “Laser Surface Hardening of Austenitic Stainless Steel”, Journal of Materials Engineering and Performance, Vol. 9, No. 5, by J.R. Ogren, Materials Park, pp.536.



19

READING THE POSSIBILITIES TO DECODE

THE MICROSTRUCTURE CHARACTERISTICS FROM MACROEXPERIENCE

Vasile MARINA *

State Technical University of Moldova,

Kishinew, Moldova Republic

ABSTRACT

The paper present the problem of practical application of relationships between the structure and proprieties. Here, it is emphasised that the symmetry elements of crystal lattice, the phases weight and the type of interatomic bond create measurable effects at macroscopic level. On this basis we succeed to formulated not only the direct problem, i.e. to deduce the constitutive equations at macroscopic level by means of constitutive equations at microscopic level, but the inverse problem, which, beside the practical importance, has a great methodological significance too. The knowledge possibilities in both directions allow us to adequately describe the material behavior as a function of external action development. The study is concentrated on the non-linear effects provoqued by the differences between the symmetry elements at macroscopic and microscopic level, prorogued by the differences between the symmetry elements at macroscopic and microscopic level, respectively, and also by the presence of several phases in the studied material, which are generated in the field of elastic strains.

KEYWORDS: Mathematical modelling, Micro-mechanics, macro-mechanics,

Structure, Proprieties

1.The principles of transition from micro-stresses and strains

to macrostresses and strains

In order to create a useful system of constitutive equations it is necessary to concomitantly study the material behavior at the level of material particle, structure element and conglomerate. We note

ijt~ , ijd~ , the stresses and strains at material particle level and tij and dij at conglomerate level; based on geometric and equilibrium equations and on homogeneity conditions at conglomerate level, we obtain the relationships [1]:

∫Δ =Δ

=0

~~1

0V ijijij tdVt

Vt ,

ijij dd ~= ; (1.1)

pqpqnmnmijij dtdtdt ==~~~~

(1..2) Beside the equations (1.1), (1.2), it is

possible to deduce another relationship, which simultaneously satisfies the geometric and equilibrium equations [2]

( ) ( )nmnmijnmnmnmijnmijij ddAddAtt −−−=−~~~~~

(1.3) where the tensor ijnmA~ depends on the material point coordinates.

If the stress ijt~ and strain ijd~ tensors are decomposed to spherical deviators and tensors:

ijijijt δσσ 0~~~ += ijijijd δεε 0

~~~+= , then

ijij σσ ~= , 00~σσ = , ijij εε ~= ,

00~εε = , nmnmijij εσεσ ~~~~ ≠ ,

0000~~~~ εσεσ ≠ (1.4)

From (1.4) it results that the macroscopic values of some physical parameters are not the same with the similar microscopic values. The difference between these parameters, named incongruity [2,3]

ijijijij εσεσ ~~~~ −=Δ

depends only on the conglomerate surface data, but on its structure too.



20

According to the proposed principle, during the real interactions, the incongruity reaches the extreme value [2,3]:

Δ

.~~~~ Extrijijijij =− εσεσ (1.5)

In expressions (1.1)- (1.5) are established direct relationships between stresses and strains at microscopic and macroscopic level. We observe that the relationships (1.1) and (1.2) do not depend on the real material structure, and (1.4) and (1.5) depend on the structure parameters. In order to obtain a correct model, we must also take into account the self-coordination phenomenon of the conglomerate deformation process. The structure elements in conglomerate loose some particular proprieties for common proprieties. As a result, it is not possible to pass directly from constitutive equations at microscopic level to macroscopic level. The self-coordination processes of thermal-mechanical interactions is emphasised by means of stresses and strains at structure elements level:

Vijij ttΔ

= ~ , Vijij dd

Δ=

~,

where VΔ is the volume of considered structure element. The relationships (1.1)- (1.3) can be written at structure elements level too. Thus, the stress and strain tensors at material particle level can be presented as follows: ijijijij tttt ~~ Δ+Δ+= , ijijij ttt −=Δ ,

ijijij ttt −=Δ ~~

ijijijij dddd ~~Δ+Δ+= ,

ijijij ddd −=Δ , ijijij ddd −=Δ~~

,

where ijij dt ~,~ ΔΔ represent the stress and strains variations in the material particles within the considered structure element ( )constdconstt ijij == , and ijij dt ΔΔ , the stresses and strains variations in the conglomerate at structure element level. We consider that between the two types of variation there is a well determined correlation: the variations within the structure element are determinate by the proprieties variation. One can observe that:

Ωijij tt = , Ωijij dd = , (1.6)

and for the variations between the stress and strain tensors we can write he relationships: t

( )nmnmijnmijij dd~A~tt~ −=− ,

( )nmnmijnmijij ddAtt −=− (1.7)

In the reversible range, ijnmijnm cA =~

,

where ijnmc represents the elasticity constants, which

are different from a particle to another within considered structure element. The fundamental relationships (1.2), (1.5) can be expressed by the variation of stress and strain tensors at the two structure levels and by the variation of respective deviators, they become:

( ) 0~~ =ΔΔ+ΔΔΩΔΩ Vijijijij dtdt ,

( ) .~~ ExtrVijijijij =ΔΔ+ΔΔ

ΩΔΩεσεσ

(1.8) One can be emphasised that for each structure element:

0~~ ≥ΔΔΔVijij dt .

As a result, the scalar product of stresses and strains variation at structure elements level is obtained with negative sign.

0≤ΔΔΩijij dt .

The integration in (1.6), (1.7) is performed according to orientation factor or to another parameter; in (1.8) was neglected the therm

( )Ω

− nmnmijnm ddA . It is important to notice that

in relationship (1.5) it is not possible to pass from the components of the deviator of stress and strain tensors at microscopic level to the deviator components of stress and strain tensors at structure element level, because

VnmVnmVijij ΔΔΔ≠ εσεσ ~~~~ .

Thus, the expression (1.5), which is an information carrier about structure element, has an extremely important physical significance. The tensor

ijnmA reflect, only those details of the interaction between the considered structure element with the other structure elements in conglomerate, which influence the material behavior at macroscopic level. The structure of tensor ijnmA and its dependence on the integration parameter (in special case, on the orientation factor ) is established from the condition that in statistical approximation the relationships (1.5)- (1.8) from a complete system of equations; on its base are deduced the constitutive equations at macroscopic level by means of constitutive equations at microscopic level.

Ω

If the relationships between stresses and strains at structure element level are know, then from (1.6)- (1.8) we can express ijt and ijd by local

thermal-mechanical characteristics, the tensor ijnmA components and the stresses and strains at macroscopic level. The problem of defining the distribution of fields ijd~ , ijt~ in the structure elements



21

is considerably simplified knowing the stress and strain average values in the structure elements, with the precision of tensor ijnmA . Thus, in relationship (1.5), beside the mechanical characteristics at microscopic level, will be present also parameters which describe the structure elements shape and size. The ijnmA components are determined from (1.5) and

(1.7). If the components of ijnmA tensor are know, then we establish the constitutive equations at macroscopic level from (1.6) - (1.8). Further on, we will refer to some examples, which demonstrate the examinated model efficiency. We will limit to the case of isotropic materials at macroscopic level, in statistical approximation, i.e. the relationships (1.8) will be written as follows: ( )( ) 0=−− ijijijij ddtt ,

( )( ) .Extrijijijij =−− εεσσ (1.9)

The material is considered isotropic at macroscopic level and the ijnmA tensor will be present as follows:

ijnmijnmijnm DAVAA 10 −= , nmijijnmV δδ31

= ,

(1.10) ijnmijnmijnm VID −=From (1.10) and (1.8) is results:

( )ijij1ijij A εεσσ −=− ;

ijij A εσ Δ−=Δ 1 (1.11)

( )00000 εεσσ −=− A ; 000 εσ Δ=Δ A (1.12) Taking into account (1.11), (1.12) and (1.7)

we obtain:

( ) 03 12

00 =ΔΔ−Δ ijijAA εεε (1.13)

2. Polycrystalline materials

with cubic lattice

The relationships (1.2), (1.3) are automatically verified for the critical cases, that is the W.Voight`s )( ijij dd = and R.Reuss`s

( )010, === AAtt ijij models. However, for the intermediate variants, the expression (1.13) imposes severe constraints not only to the ratio A0/A1 , but to the structure of relationships between the stress and strain variations too.

Because ( ) ,0,020 ≥ΔΔ≥Δ ijij εεε it

results that 01

0 >AA

and consequently, the

E.Kroner`s model:

,5457

1 νν

−−

= GA GA 40 −=

is in a qualitative contraction with the fundamental relationship (1.13). We can observe that for the polycrystallines monophasic materials with cubic lattice ( 00 εσ K= , 00 εσ K= ), the spherical tensor variations cannot be determined based on the relationship (1.12), which is reduced to the expression: ( )( ) 0000 =−− εεAK (2.1) Indeed, from (1.13) it results that for

00 εε = we obtain ijij εε = , i.e. the Voight`s model. A complete system of equations, supposing that

ijij tt ≠ , ijij dd = , is obtained only if it is assumed that in (2.1) and (1.13):

,0 KA = ( ) ( )( ijijijijKA εεεεεε −−=−

312

00 )

(2.2) Thus, for some materials, the relationships

between the stress and strain variations are determined based on the expressions (1.11) and (1.12). In the case of polycrystalline polyphasic materials with cubic lattice, from (1.12) we obtain:

,00

000 εεεε fff

AKKK

−

−=−=Δ

(2.3)

KK f ≠

where Kf is the compression modulus in the phase no. “f”, and fA0 is the internal parameter, which characterises the spherical tensor field unhomogeneity inside the considered phase. For the spherical tensor variations, from (1.13) and relationships:

0εσ ff K= , 00 εσ K= , it is obtained the expression: ( )( ) ( )( )ij

fijij

fij

ff AKK εεεεεεεε −−=−− 100003 (2.4)

If ., than from (1.13) and (2.3) we obtain the following relationship:

constKK f =≠≠ ,00ε

( )( )ijf

ijijf

ijf

ff

ff A

AKKKA εεεεε −−=⎟⎟

⎠

⎞⎜⎜⎝

⎛−−

120

2

003

(2.5) from which it results that the parameter fA0 varies from a particle to another and also within the each phase in the case of polycrystalline polyphasic materials. In the relationships (2.4) there are only invariant parameters; because of this, it is appropriate to calculate the scalar product of the variations of strain tensor deviator within the crystallographic system of



22

coordinates x`i for each structure element. For the polycrystalline materials with cubic lattice, within the

crystallographic system of coordinates it is obtained the following expression:

( )( ) ( ) ( ) ( )[ ]+′+′+′⎟⎟⎠

⎞⎜⎜⎝

⎛+−+−

=′−′′−′ 222

222

211

2

11211

12112 εεεεεεε fff

ff

ijijijij AccccG

( ) ( ) ( )[ ]2

232

132

12

2

144

44

2222 εεε ′+′+′⎟⎟

⎠

⎞⎜⎜⎝

⎛+−

+ ff

f

AccG

(2.6)

To write the expressions (2.6) there were

taken into account the relationships between stresses and strains within the crystallographic system of coordinates:

( ) ,11121111 εσ ′−=′ cc …; ,2 124412 εσ ′=′ c (2.7) After the substitution of these expressions and formula ijij Gεσ ′=′ 2 in (1.11), we obtain:

( ) ,2

11211

11111 Acc

GA+−′+

=′εε

( ) ,2

2

144

2112 Ac

GA+

′+=′

εε (2.8)

Within the overall system of coordinates

nmjminij aa εε =′ and, as a result, the relationship (2.6) get the form:

( )( )

qknmji

jkiqjminff

f

qknmi

ikiqiminfff

ff

ijijijij

aaaaAccG

aaaaAccccG

εε

εεεεεε

⎟⎟⎠

⎞⎜⎜⎝

⎛⎟⎟⎠

⎞⎜⎜⎝

⎛+−

+

+⎟⎠

⎞⎜⎝

⎛⎟⎟⎠

⎞⎜⎜⎝

⎛+−+−

=′−′′−′

∑

∑

=≠

=

3

1

2

144

44

3

1

2

11211

1211

222

2

(2.9)

From (2.9) it results that the scalar product

of the variations of strain tensor deviator depends on: the crystals elasticity constants , which varie form a phase to another; the internal parameter A

fff ccc 441211 ,,

fl (reflects the unhomogeneity of ijε field in the

conglomerate); it is worth too considering the variant when or 11 AA f = ( )f

ff AfAA ,11 ψ= (where Af

is the anisotropy coefficient of “f” phase and fψ the phase weight); the crystallographic axes orientation ( )jiji a x,xcos,`x i′= ; the values of the

components of strain tensor deviator ijε (at macroscopic level). In the considered phase crystals, the variations of the strain tensor deviator change from a crystal to another only function of crystallographic axes orientation. In (2.9), G and Af

l are unknown parameters. To calculate them, we`ll use the relationships (1.6), (1.9), (1.11), (2.8) and (2.9). We can observe that the relationships (2.8) within the overall system of coordinates could be written like an unique expression:

kqnm

mqnkmjniff

f

nnqnknjnifff

f

ij aaaaAcGAaaaa

AccGA εε ⎥

⎦

⎤⎢⎣

⎡++

++−

+= ∑∑

=≠=

3

1144

13

111211

1

222

(2.10)

Substituting (2.10) in (1.6) and taking into account the relationships:

( )jkiqjqikkqijnqnknjni aaaa δδδδδδ ++=Ω 15

1

, (2.11) we will obtain:

ikjqjqkikmqmjnkni Iaaaa151

152

−=Ω

δδ , (2.12)

where fψ is the “f” phase weight and N-phases number. To deduce the formula (2.13), it was considered that the material particles distribution is homogeneous from a statistical point of view, during each phase. Substituting (1.11) in (1.9) and taking into account the relationship (2.9) in the obtained expression, after integration, we find:



23

ExtrGAc

cGAccccGAN

ffff

f

fff

fff

=⎪⎭

⎪⎬⎫

⎪⎩

⎪⎨⎧

⎥⎥⎦

⎤

⎢⎢⎣

⎡⎟⎟⎠

⎞⎜⎜⎝

⎛+−

+⎟⎟⎠

⎞⎜⎜⎝

⎛+−+−∑

= 21

222322

51

2

144

44

2

11211

12111 ψ (2.14)

Based on the relationships (2.13) and (2.14),

we determine the unknown constants 2G and

values. If is known, then from the relationships

(2.5) and (2.9) we calculate the parameter

fA1fA1

fA0 values and from (2.3) after integration:

00

=−

−ff

f

AKKK

(2.15)

we establish the compression modulus K values. We can observe that form (2.5) and (2.9) it results that

fA0 depends not only on the crystal’s elasticity constants, but on the deformation degree too. Consequently, we determine that the compression modulus at macroscopic level will depend on the

deformation degree in the case of polyphonic materials with cubic lattice.

If we refer only to monophasyc polycrystalline materials, then form (2.13), (2.14), they result the relationships:

( ) ( )[ ]121144

1211441211441 3

34ccc

ccccccA−+

−+−= ,

(2.16)

( )( )( ) RV GG

cccccccccG =

−+−+−

=121144

121144121144

343

(2.17)

( )( )

2

2332516 ⎟

⎟⎠

⎞⎜⎜⎝

⎛

+++−

−=−

AAAA

nmnm

ijijijij

εσ

εσεσ, (2.18)

1211

442cc

cA−

= , ( )( ) ijijAAA

A εσεσεσ2

0000 2332512 ⎟

⎟⎠

⎞⎜⎜⎝

⎛

+++−

=− , (2.19)

obtained for the first time in [4]. We mention that Gv and GR represent the shear modulus values, obtained by W.Voight (1928) and A.Ress (1929). The relationship (2.19) allow us to evaluate the influence of anisotropy factor A on the energy dissipation under cyclic loads in elastic field. Based on the relationships (2.16), (2.17) and the formula for anisotropy factor, we establish the following relationships:

6332

21

++

=AA

GA

(2.20)

from which it results that the internal parameter Af , which emphasises the inhomogeneities of strain and stress deviator fields in conglomerate satisfies the inequality

GAG34

1 << (2.21)

If the microscopic characteristics are unknown, then one can assume that . This type of relationships are useful for numerical computations, performed based on the expressions (2.13) and (2.14), in the case of polycrystalline polyphasic materials.

GA 1,11 ≈

We observe that the A1 values, obtained within the studied model are always smaller than A1

k values, which result from E.Kroner`s model 5.00÷=ν .

( ) GAG k 375.1 1 << (2.22) According to expression (2.18), for N=1 it is obtained the relationship:

( ) ( )( ) 1121144

1441211121144

53435Accc

AccccccG+−++−+−

=

(2.23) from which it results that the shear modulus diminishes with A1 increasing; as a result, the shear modulus obtained from (2.17) is “more microscopic” than the shear modulus determined by E.Kroner`s model. We mention that the A0 values are different from qualitative point of view within the 2 models:

( ) .4; 00 GAKA k −==

3. Polycrystalline materials with crystalline lattice symmetry lower

than the cubic one For the polycrystalline materials with crystalline lattice symmetry lower than the cubic one, the variation of spherical tensors 00 ,εσ is obtained based on the relationships (1.12). As a result, the formula (1.13) is completely verified.



24

According to [3], the symmetry elements at microscopic level lead to the appearance of some effects at macroscopic level, based on which one can concludes about material structure. Thus, from studied model it results that each structure element at microscopic level determines effects measurable at macroscopic level. Based on this kind of effects, it becomes real the inverse problem: the decoding of

some details about material microstructure from macroexperience. We mention that the inverse problem was considered unreal until establishing these effects. In order formulate the final conclusions, we`ll refer to the computation relationships obtained in [3] for the polycrystalline monophasic materials with crystalline lattice symmetry lower than the cubic one:

( ) 220

22

1

020 6036013 bbbbb

bb

AA

bbb pqllpqkkpqklpqklijmmijnn −−=⎟⎟⎠

⎞⎜⎜⎝

⎛−+−

μ (3.1)

the bijnm tensor is determined based on the relationships:

( ) ExtrAbAbbbb ijmmijnn =

−

−0

1

20

213

, (3.2)

where

,31

0 kknnbb = ( )kknnknkn bbb −= 3301

, pqpqεε

εμ

202 = (3.3)

The bijnm tensor is determined based on the relationships: ( ) ijnmmkmkijkl IAcb =− lnln , ijnmijnmijnm DAVAA 10 −= (3.4)

In the system (3.4) cklnm are elasticity

constants at microscopic level within the crystalographic system of coordinates. Thus, in the non-linear equation system (3.1), (3.2) are present two unknowns, A0 and A1; based on them, from the relationships:

,31

0 kknnbb = (3.5)

( ) ,122 1 =+ AGb (3.6) we calculate the macro-elasticity constants K and G. From (2.6), (2.7) it results that A0 and A1 depend not only on the microscopic elasticity constants, but on the deformation degree too. As a result, we establish that the relationships between the stresses and stains at macroscopic level are non-linear, if the structure elements have a lower symmetry than the cubic one; the elasticity characteristics K and G at macroscopic level depend

on the deformation degree pqpqεε

εμ

202 = or stress.

4. Conclusions

From the structural model it results that each structure element at microscopic level determines a measurable effect at macroscopic level. Starting from the established relationships between cause and effect, based on the experience at macroscopic level, we succeed to formulate conclusions about the microstructure of the examined material. If from macro-experience we establish that the shear (G) and compression (K) moduli are not influenced

by the stress or deformation degree, than we can assert that the examined polycrystalline material is monophasic with cubic crystalline lattice. In the case of polycrystalline polyphasic material with cubic lattice, from macro-experience we find that the shear modulus is not influenced by the stress degree and the compression modulus K depends on the stress/strain degree. If from macro-experience we establish that both elasticity characteristics (K,G) depend on the stress degree, from the proposed model it results that the crystalline lattice symmetry of the examined polycrystalline material is lower than the cubic one; determining the structure of K and G parameters dependence on stress/strain degree, one can establish the details about the symmetry elements of the crystalline lattice and the presence of one or more phases. We mention that the listed effects do not result from other models of transition from the microscopic state to the macroscopic one. Based on the established non-linear effects, one can assert that exists an equivalence between the direct problem, i.e. the deduction of constitutive equations at microscopic level, and the inverse problem, the decoding of thermomechanical characteristics at microscopic level from macro-experience. We can observe that the majority of specialists, until the establishing of mentioned non-linear effects, consider the inverse problem without solution. The existence of inverse solution has a great practical and scientific importance. Because the structure elements in conglomerate modifie some proprieties, the knowledge direction from micro-to macro especially during the irreversible processes, leads to some unavoidable errors in conglomerate behavior description. This finding is available in the case of



25

inverse problem too, because not all the details of structure elements can be precised from macro-experience. Thus, the knowledge process will become more complete, if the study is performed in both directions. We mention that from macro-experience it can be determined the interatomic bonds type. Thus, in the case of ionic bond (the model of central interaction proposed by Cauchy), besides the symmetry relationships, there are deduced the following expressions [5] too:

;,, 232322331313113312121122 cccccc === (4.1)

;,, 332123312213122311321312 cccccc ===

Thus, determining, e.g., from macro - experience on a NaCl or KCL crystal that c1122=c1212, (for the materials with cubic lattice the 6 relationships (4.1) are reduced to an unique one c11=c44, 1∼11, 4 ∼12) , we establish that the interatomic bond of examined material is an ionic one. For the materials with other

bonds type, are not deduced yet supplementary relationships (4.1) type: this problem has a great methodologic importance in the context of stated conclusions.

References [1]. Hill R. The Elastik Behavior of Cristalline Aggregate, Proccedings of Phisical Society. London A, 65, p.349. 1952. [2]. Marina V. Stat Equation under Proportional Nonnisothermal Loading, Internation Applied Mechanics. Vol.33, Nr.2, 1996, p.92-99. [3]. Marina V. The principles of passing From the Microscopic state to the Macroscopic one, Metallurgy and New Materials Researches. Vol.IV, Nr.4/1998, p.14-28. [4]. Marina V. Prinţipî perehoda ot micro c macro napreajeonno-deformirovannomu sostoianiu. ` Buletinul Academiei de ştiinţă a R. Moldova. Matematica. Nr.2, 1998, p.16-24. [5]. M.L.Bernştein, V.A.Zaimovskii. Mehaniceskie svoistva materialov. Moscova: Metalurgia, 1979,495p.



26

SIMULATION OF SEMI-PLANETARY ROLLING PROCESS

AT EXPERIMENTAL STAND

Nicolae CĂNĂNĂU, Alexandru IVĂNESCU, Lilica IVĂNESCU

“Dunarea de Jos” University of Galati, 47 Domneasca, Romania email: [email protected]

ABSTRACT

The semi-planetary rolling is a original method for plastic deformation of the

semi-product. The semi-planetary rolling mill consists of a planetary rolling cylinder (superior cylinder and a massive cylinder (inferior cylinder). In the proximity of the planetary cylinder, the strain intensity is grater then the strain intensity developed in the proximity of the massive cylinder. Consequently into the deformed material an asymmetrical deformation state is developed. The semiproduct curves and the curvature is influenced by the thickness reduction in the semip-planetary rolling process and b/h ratio. In this paper are showed the experimental researches concerning the influence of the named parameters on local deformations and curvature of the semi-planetary rolled product.

KEYWORDS: semi-planetary rolling, asymmetrical deformation

1. Introduction

The semi-planetary rolling, (Figure 1) leads to an asymmetrical stress and strain states on the thickness of the semi-product.

The strain developed in the proximity of the superior cylinder (planetary cylinder) have greater intensity then the strains developed in the proximity of inferior cylinder (massive cylinder). As effect, after the semi-planetary rolling the initial semi-product, with rectilinear form, becomes curved. And the initial body, with curve form, becomes rectilinear or curve body, the last with different curvature, after the semi-planetary rolling.

For example an application of semi-planetary rolling, which may become an important solution for quality assurance of the metallurgical products, may be the unbending of the continuous cast slabs of high strength steels. These steels have a reduced plasticity, in casting state and, also, at the temperature of 850 – 950 0C, they present a fragility domain.

At the same time in the unbending zone, into the proximity of interior surface of slab section (towards the curvature point), a plane tensile stress state is developed.

Coupled these two causes, respectively, the reduced plasticity of steel and the bi-axial tensile stress state may lead at the appearance of the superficial fissures.

The semi-planetary rolling applied in the unbending zone, to the exit of slab from curve thread of the continuous casting machine, induces an

unbending moment in the condition of the compression stress state.

R

lc′

R1 2

1d

N

3

4

v0h0 h1

Fig. 1 Scheme of semi-planetary rolling:1-suport planetary cylinder, 2-satellyte cylinder,3-massive cylinder, 4-semi-product

In conditions of a good correlation of the

thickness reduction, according to the radius of curvature (or curvature) of continuous casting



27

machine and ratio of breadth and thickness, the slab may be unbend through the semi-planetary rolling directly. From the forces equilibrium (see Figure 2), results:

mp FFF == (1)

Fig. 2 Forces scheme of semiplanetary rolling

Fp

Fm

F

zp

zm

h

d

D

am

ap

Fp

Fm

In relation (1) Fp is the force developed on planetary cylinder and Fm is the force developed at the massive cylinder. It is evident :

bapF

bapF

mmm

ppp

⋅⋅=

⋅⋅= (2)

and p

m

m

paa

pp

= (3)

It denominate zp the penetration depth of the

satellite cylinder and zm the penetration depth of the massive cylinder into material of bodi. From geometrical conditions we have:

and

2mmm

2ppp

zzR22a

zzr22a

−⋅=

−⋅=

(4)

In the relations (4) r is the radius of satellite

cylinder and R is the radius of the inferior cylinder. Because the terms zp

2 and zm2 are very

small then the products r·zp and R·zm , respectively. Thus we can neglect the second terms of the

radical arguments and the relations (4) become: mp zR22ma,zr22pa ⋅=⋅= (5)

From (3) and (5) we obtain:

p

m

m

pzrzR

pp

⋅⋅

= (6)

Using the relations (5) we can write:

m

p

m

pzRzr

aa

⋅

⋅=

From this relation we obtain:

2

m

p

m

paa

Rr

zz

⎟⎟⎠

⎞⎜⎜⎝

⎛⋅= (7)

.

zp

zm

εhp

εlp

εhm

εlm

Fig. 3 The local deformation states

Consequently, in the cross section of the

semi-product will be developed the compression stress with greater intensity in the proximity of planetary cylinder.

Thus into the cross section of the body a bending moment is developed (Figure 4). And intensity of this moment is greater, so much greater is the reduction of thickness (Δh = h0 – h1, h0 is the initial and h1 final thickness.).

At the same time, the bending moment is influenced by b/h ratio. For the great b/h ratio, the condition of the plan deformation state is performed and the relations (8) are satisfyed.

For smaller b/h ratio, relatively, the deformation in the breadth direction are different of zero and the bending moment intensity is more reduced.



28

Fig. 4 Bending process at semi-planetary rolling

Thus the intensity of the bending moment is influenced by reduction of thickness Δh and b/h ratio:

⎟⎠⎞

⎜⎝⎛Δ=

hb,hfMi (10)

If the intensity of the bending moment is

great the rolled product curves more intense. The bending intensity of the rolled product

may be defined by the curvature or radius of curvature:

⎟⎠⎞

⎜⎝⎛Δ==

hb,hg

R1C

ii (11)

where the Ci is curvature and Ri is radius of curvature of semi-product rolled of semi-planetary rolling.

In this paper are show the experimental results concerning the influence of the thickness reduction and the b/h ratio on curvature to semiplanetary rolling.

2. Experimental conditions

The experiments were been done at a laboratory rolling mill equipped with upper planetary cylinder. The exterior diameter of the planetary cylinder is of 80mm. The diameter of the sarellite cylinders is of 12mm. The planetary cylinder includes 18 satellite cylinders.

The active surface of the inferior cylinder (massive cylinder) is circular and the diameter of this surface is of 80mm. The length of the active surface of cilinders is of 50mm.

The samples were processed from lead, through casting and symetrical rolling process at the dimensions: the thickness of 10mm and various values for breadth. The values of the b/h ratio were: 1.15, 2.5 and 3.3.

The values of the thickness reduction were of 0.2, 0.4 and 0.6mm i.e. 2%, 4%, 6%, respectively.

Other matter of interest for experimental researches was thw repartition of penetration depth of the sarellite cylinder and massive cylinder.

With this end in view, it was prepared an equipment compused of a massive (inferior) cylinder and a small cylinder, with the diameter of 12 mm, and a sensitive tensometer. Also, we prepared samples with different values for b/h ratio. Loading the experimental system to the different press forces shall been induced various penetration depths of upper and lower cylinders according to their diameters.

3. Experimental results and interpretation

In the aim of evaluation of the zp/zm ratio they

were effectuated pressing tests using the samples of lead between the massive cylinder with 80 mm of diameter and satellite cylinder with 12 mm of diameter. The Table 1 contents some results of these experimental researches.

The values of the penetration depths were been calculated using the geometrical relations: It is evident that the values of zp/zm ratio are 1.72 – 2.18. This proves that the penetration depth in the proximity of the planetary cylinder is greater of the penatration depth in the proximity of the inferior massive cylinder.

2

m2m

2p2

p

2aRRz

2a

rrz

⎟⎠⎞

⎜⎝⎛−−=

⎟⎟⎠

⎞⎜⎜⎝

⎛−−=

(12)

Table 1

b0 mm

h0 mm

ap mm

am mm

zp mm

zm mm m

p

zz

24.0 9.2 3.2 6.2 0.22 0.12 1.83 24.0 9.2 4.7 9.3 0.48 0.22 2.18 33.5 6.7 3.0 5.9 0.19 0.11 1.72 33.5 6.7 4.5 8.6 0.44 0.24 1.80

Consequenly the effect of the asymmetrical

strains repartition at semi-planetary rolling is justifyed.

Also we effectuated experimental tests in the aim of establishing the influence of the thickness reduction and b/h ratio on curvature and radius of curvature of the semi-planetary rolled samples. In Figure 5 is showed the radius of curvature variation of the samples rolled at the semi-planetary rolling mill in function of thickness reduction and b/h ratio.



29

Fig. 5 Radius curvature variation withΔh and b/h ratio

0

20

40

60

80

100

120

140

160

180

200

0,022 0,040 0,056 0,075 0,084

Specific thickness reduction, Dh/ho

CurvatureRadius,[mm]

6810

35455565758595

105115125135

0 0,02 0,04 0,06 0,08 0,1

Reducerea relativa, delta h/ h0

Razadecurbura,R,[mm]

Fig.6 The overall variation of the radius of curvature with reduction of thickness

5

10

15

20

25

30

0 0,02 0,04 0,06 0,08 0,1

delta h /h0

1/R

Fig. 7 The overall variation of the curvature with the reduction of thickness

In Figure 6 is showed the cumulated radius of curvature variation of the samples rolled at the semi-planetary rolling mill in function of thickness reduction.

In Figure 7 is showed the cumulated curvature variation in function of thickness reduction.



30

In Figure 8 is showed the curvature

variation of the samples rolled at the semi-planetary rolling mill in function of b/h ratio for various thickness reductions.

The curvature of the samples increases when the thickness reduction increases and, also, when the b/h ratio increases.

The influence of the reduction of thickness has a more important influence then the b/h ratio.In case of the semi-product, as the thick sheets or continuous casted slabs, when the b/h ratio has great values, the influence of the b/h ratio is favorable for the productivity of the semi-planetary rolling process. Haw the effect of the stress and strain fields, developed at the semi-planetary

rolling, is reversible this deformation method may be used, also, for the unbending of the continuous casting slabs of high strength steels, without tensile stresses and, consequently, in conditions of the guarranted quality of production. Thus, the quality of continuous cast slabs shall be assured, because the unbending process is developed in same time of semi-planetary rolling, in conditions of compression stresses state. Since the critical tensile stresses are eliminated, the cause of the superficial fissures is eliminated and the surfaces of the continuous cast slabs results continuos without fissures or cracks.

0

5

10

15

20

25

0 1 2 3 4 5

b0/h0

1/R

Fig. 8 The variation of curvature with the b/h ratio

4. Conclusions

The semi-planetary rolling method may

be applied to processing of the curve pieces or for unbending of the curved semi-product. For example the unbending of the continuous casted slabs of the high strength steels may be a favorized process of semi-planetary rolling.

The researches showed in this paper prove the validity of idea of continuous tast slabs unbending.

The advantage of this application is the unbending in condition of the compression stress state, this being the condition of the quality assurance of the casted slabs and, finally, the guarrantee of the metallurgical products quality.

In case of curved pieces production, for the deffined curvature (or radius of curvature) and b/h ratio we may be establish the optimum value of the thickness reduction.

In case of the unbendig of continuous casted slab (with the b/h ratio and radius of curavture deffined) we may be establish the thickness reduction necessary for its unbending.

The researches performed and showed in this paper prove that the semi-planetary rolling is a viable method for these aims.

References [1]. Cănănău,N.,Ivănescu,A. Studii şi cercetări cu privire la defectul discontinuităţi la tablele groase din oţel. Sesiunea de comunicări “Realizări şi perspective în Metalurgie şi Ştiinţa materialelor”. Universitatea din Galaţi, 13-15 octombrie 2000 [2]. Cănănău,N.,Ivănescu,A. Study on the Relation between the Quality of Continuous Cast Slabs and the Surface Discontinuities of the Steel Thick Sheet. ’01 International Conference on Advances in Materials and Processing Technologies, September 18-21, 2001, Leganes, Madrid, Spain, vol. 2, p.949-954. [3]. Cănănău,N.,Ivănescu,A. Technological Method of Forestall of the Superficial Cracks Network to the Continuous Cast Slabs.. ’01 International Conference on Advances in Materials and Processing Technologies, September 18-21, 2001, Leganes, Madrid, Spain, vol. 2, p.955-960. [4] Anon, Thin Hot Strip Using Planetry Mill, Steel Time Intern., Sept. 1995, 33-5. [5] Cananau,N.,Ivanescu,A.,Cananau,D., Study on the Deformation Process at the Semi-planetary Rolling, Acta Mechanica Slovaca, Kosice, nr.3, 2001.



31

EXPERIMENTAL RESEARCHES ON THE Zn – Fe COATING

Tamara RADU, Olga Mitoşeriu Lucica BALINT

“Dunarea de Jos” University of Galati, 47 Domneasca, Romania

email: [email protected]

ABSTRACT

The protection alloyed Zn-Fe layers are recommended for the following qualities: a better weldability than the zinc-coated plates, very good capacities for painting and lacquer wear resistance, good adherence, lower zinc consumption by layer thickness below 100g/m2. This paper intent to evaluate the corrosion resistance of these coatings as compared with the zinc-plated steels a basic criterion for the assessment of a metal coating. For the protection layers made from Zn - Fe alloys an improved corrosion resistance is found when the amount of Fe is increased along with a higher corrosion resistance as compared with the zinc-coated plates and microstructure, adherence, thickness layer, Zinc-coated and Zn-Fe alloy samples were testes by two procedures:

a) Electrodepositing b) Heating the zinc-coated plate with the Fe diffusion into the protection layer. The corrosion resistance of the samples was tested by two procedures: the test in salty spray and the electrochemical method:

KEYWORDS: Zn-Fe layers, zinc-coated plates, chroming, phosphating, salt spray, Tafel law.

1. Introduction

About 20% of the steel products in the world are zinc plated and there are good predictions as regards the future of zinc plating [1]. The increasing demand of zinc – coated plates instead of cold rolled plates led to an increased production of good quality zinc depositions.

Zinc improvement has been achieved by alloying it with one or more elements (Al, Mg, Ni, Mn, Co, etc) or by subsequently treating the zinc-plated surface (heat treatment, passivation or organic surfaces coating) [2].

Diversifying the range of zinc-plated products resulted in: - Lower reactivity of the sacrifice metal (Zn) by substituting it with a less electronegative material, which should corrode slower or provide longer protection at an identical or smaller layer thickness. - Improvement of a property such as weldability, painting ability, temperature resistance, layer plasticity, etc.

The protection alloyed Zn-Fe layers meet a large part of these requirements and are recommended for the following qualities: a better weldability than the zinc-coated plates, very good capacities for painting and lacquer wear resistance, good adherence, lower zinc consumption by layer thickness below 100 g/m2.

This paper intent to evaluate the corrosion resistance of these coatings as compared with the zinc-plated steels a basic criterion for the assessment of a metal coating.

2. Experimental Researches

Zinc-coated and Zn-Fe alloy samples were testes by two procedures: a) Electrodepositing b) Heating the zinc-coated plate with the Fe diffusion into the protection layer. a).To prepare the electrodeposition of the Zn-Fe alloys, use was made of electrolytes of simple salts based on chlorides and sulphides with allowances to increase the electrolyte conductivity, screening the Fe3+ ions and buffer substances. The solutions used and the working conditions are presented in Table 1. Layers of Fe content ranging between 10% and 40% were obtained. b).The zinc-coated plate was heated to 550-650 OC exposing it for 1-3 minutes, when Zn - Fe alloy layers of 11-22% Fe was obtained [3]. The researches were also focused on the identification of some passivation procedures aimed at increasing the of the Fe Zn alloys obtained by diffusion while improving their aspect as well. The passivation solutions and the application conditions used for this purpose are given in Table 2 and Table 3.




32

Table 1-Electrolysis solution and Zn Fe alloys electrodeposition conditions: Nr. crt.

Electrolyte Concentration Condition

1 FeCl2 + ZnCl2CH3COONa Acid citric

300 15 – 35 5 - 10

pH = 1,5 – 2,5 t = 20 – 50 0C

I = 10 – 30 A/dm2

2 FeSO4 + ZnSO4Na2SO4

CH3COONa Acid citric

500 30 – 50 15 – 20 5 - 10

pH = 3

t = 50 0C I = 30 A/dm2

3 FeSO4 + ZnSO4 (ZnCl2) MgSO4Na2SO4

Acid citric

500 30 – 50 40 – 50

5

pH = 2,0 t = 20 – 50 0C

I = 10 – 30 A/dm2

Table 2 Passivation solutions on Zn-Fe alloy coated –plates

Solution 1 Chroming K2Cr2O = 1 g/l NaF = 5,0 – 5,5 g/l

HNO3 conc. = 8,0 – 15,0 cm3/l Solution 2 Chroming Na2Cr2O = 1 g/l

H2SO4 = 6,0 – 8,0 cm3/l Solution 3 Phosphating H3PO4 = 20 - 30 g/l

ZnO = 20 - 25 g/l HNO3 conc. = 1,5 – 2,0 cm3/l

Table 3 Condition of passivation on Zn-Fe alloy coated –plates Solution Temperature,[ 0C] Time, [min] pH Colour

Solution 1 18 - 25 0,5 1 - 2 Blue Solution 2 18 - 25 0 1 – 2 1 – 2 Green Solution 3 28 - 30 25 - 30 2,4 – 2,8 Black

The corrosion resistance of the samples was tested by two procedures: the test in salty spray and the electrochemical method: • The salt spray test:

The samples were introduced in a corrosive environment at 35±2 oC and with a relative humidity for 460 hours.

After each 24-hour cycle, the samples were washed, cleared of any corrosion products and weighed. The corrosive atmosphere was obtained by the fine pulverization, in the Kesternich equipment, of a solution of composition: NaCl-27 g/l, MgCl2 -6g/l, CaCl2 -1g/l, KCl -1g/l pH = 6.5÷7.2. • The electrochemical method

For this method was used a TACUSSEL apparatus and an X-Y PHILIPS registering in the following conditions: H2SO4 0.1 N solution, the polarization velocity is 250 mV/min, reference electrode ESC, a platinum counter-electrode and the temperature of 20 oC.

The polarization was made starting from values that are more negative than the corrosion potential (-1050 mV/ESC). The potential varied with constant speed/ velocity and corrosion current was registered.

3. Results and Conclusions • The salt spray test:

The corrosion velocity was expressed by the gravimetric index, i.e. the sample mass variation as a result of corrosion with respect to the surface and time unit. The results obtained are graphically illustrated in Figs.1-5.

For the protection layers made from Zn Fe alloys obtained by heating the Zn-coated plates an improved corrosion resistance is found when the amount of Fe is increased (Fig 1) along with a higher corrosion resistance as compared with the zinc-coated plates.

In the case of electrolytic depositions, where the layer thickness is lower than in the previous cases and the Fe is higher, the corrosion resistance decreases for amounts of Fe higher than 20%; as a whole the corrosion resistance is still higher than that of the heat Zn coated plates (Fig 2)



33

Fig.1. Corrosion behaviour of the Zn and Zn-Fe alloy coated samples by heating the Zn coated plate

Fig.2. Corrosion behaviour of Zn and Zn Fe alloy coated samples, made during electrodepositing

Passivation, in the tested solutions, increases the corrosion resistance as compared with the unpassivated samples when the coated samples are cleaning by oxalic acid in chroming solutions (solutions 1, 2) and tartric acid in phosphating solution (solution 3).

• Electro – chemical test The electro-chemical measurements were

made to check the results from the salt spray test. In order to assess the corrosion resistance the anode polarization curves were plotted and the electro-chemical parameters are given in Table 4.



34

T O T T TO

Fig.3. Corrosion behaviour of Zn and Zn Fe alloy coated samples, passivated in solution 1.

T O T T TO O O

Fig.4. Corrosion behaviour of Zn and Zn Fe alloy coated samples, passivated in solution 2.

The Tafel slope (b) is determined by the Tafel law: η = E-E0 = b (log i – log i0) (1) b = E-E0 / (log i – log i0) (2)

The angle tangent of the linear part of each curve was calculated.



35

T O T TO

Fig.5. Corrosion behaviour of Zn and Zn Fe alloy coated samples, passivated in solution3.

Table 4. Electro-chemical parameters:

Layer composition Corrosion potential, εc, [mV/ESC] Tafel slope, b, [mV] Zn (4,3%) Fe -770 516

Zn-Fe (11,25%) Fe -700 733 Zn-Fe (15,72%) Fe -700 733 Zn-Fe (22,28%) Fe -660 332

From the polarization curve analysis, it results that: - for all samples a general corrosion process becomes obvious at the anode polarization as shown by the high corrosion current values; - At the Zn Fe coated samples the corrosion potential tends to high electrically positive values which suggests a better corrosion behaviour as compared with the zinc coated plate as shown by the salt spray test too; - Although indirectly, the small values of the Tafel slope for the Zn-Fe (22,28%Fe) samples indicate a slower corrosion kinetic than with the other samples; - From the over tension variation with respect to the current intensity the following order is established:

η1 < η2 ≈ η3 < 1 which proves a higher corrosion resistance according to the same order.

The microscopic analysis of the samples measured electrochemically (anode polarization) revealed different aspects: • heavy corrosion on the Zn samples • in addition to general corrosion, a very marked net-shaped attack on the alloy Zn- Fe samples for 22,28% Fe.

As indicated by the investigations made, the corrosion resistance increases for the Zn -Fe alloy

layers as compared with the Zn samples depending o the amount of Fe in the surface layer.

During the corrosion process the zinc forms a semi-conducting oxide of type “n” with cathions in excess. The Zn Fe alloying (as in the case of Zn Fe depositions) results in an excess of positive charges. Substituting two ions of Zn2+ from the crystalline network of the Zn oxide with two ions of Fe3+ results in two positive exceeding charges which requires, for compensation purpose, that a Zn2+ ion leaves the spaces where it is in excess.

This leads to a diminished concentration of the net defects and implicitly to a lower corrosion velocity. Even if a higher concentration of Fe leads to improved corrosion resistance, it must be taken into account that there are other properties negatively affected. Thus a significant increase of the layer hardness and implicitly of the pulverising index is obvious [4] thus limiting the use of these products to hose having 8-14% Fe in the layer.

References [1].C.CabrillaC, P. Ponthiaux, LesEntretiens de la Tehnologie, vol.I, Paris, 15-16 Martie 1994. p 3-21. [2]. R. Muller, Galvano Organo, 649, (oct 1994), p 727-731. [3]. T.Radu, L. Benea, Euromat'98, Lisabona (1998) p 509 – 515. [4]. D.C. Bae, S. Choi, C. Shinj, Galvatech'92, Amsterdam (1992), p 112-116.



36

RESEARCHES CONCERNING

HEAT TREATMENT EFFECTS UPON THE MORPHOLOGY OF HIGH-TENSILE FOUNDRY BRASS STRUCTURE

Adriana PREDA1, Dionisie BOJIN 2, Sanda Maria LEVCOVICI3

1 C.S.UZINSIDER Engineering A/S Galati, Politechnical Institute, Bucuresti, 2

3 “Dunarea de Jos”, University, Galati E-mail: [email protected]

ABSTRACT

This paper aimed to emphasize the structural modifications foundry brasses undergo

as a result of diffusion and oversaturation as well as secondary hardening subsequently to practical application of several heat treatment conditions. Experimentally, at distinct temperature values, several hardening and annealing were applied on three brasses alloyed with Al, Fe, Mn, single-phase β and two-phases α+β with different rates of α phase, in order to analyze their microstructure, hardening and distribution of alloying constituents. The results which have been obtained further to the researches carried-out with the aid of an electronic microprobe have enabled us to determine heat treatment response susceptibility of high-tensile foundry brasses as well as the effects upon their structural morphology.

KEYWORDS: Foundry brasses, heat treatment, microstructure

1. Theoretical considerations Transformations which took place in solid over-

saturated solutions (putting in solution followed by precipitation) are the basis for alterations of chemical composition, structural morphology and properties corresponding to all heat treated non-ferrous alloys [1, 2, 3].

Selection and application of several heat treat-ment conditions to brasses are based on aspects related to hammer hardening of solid solutions [4, 5, 6], through both combined effect of alloying constituents as well as precipitation of α phase particles according to the following pattern:

dissolution of alloying constituents in crystal lattice leads to a hammer hardening of solid solutions by slowing-down the dislocation movement showing conducive effects to strength and hardness;

homogeneous repartition of α phase particles distributed as fine needle-like shape, as well as inter and intra-granular lamellas, generate the same hammer hardening of solid solutions by alteration of internal stress condition leading to barriers built on the way of dislocations.

Hardening achieved by precipitation is one made by dispersion, and is justified by interaction between dislocations and precipitation particles.

It results from slowing-down dislocations movement caused by internal stress fields defined by secondary phase particles, these acting as obstacles on the way of dislocation movement.

Distribution’s degree together with particles shape play an important role in hardening action of precipitated particles. Sphere-like shaped particles have shown the weakest effect of hardening.

Reference literature [7, 8, 9] emphasized the fact that heat treatments are especially applied to deform-ing brasses in order to obtain adequate structures, eliminating the anisotropy of mechanical properties.

For high-tensile foundry brasses, these heat treatment conditions have more likely a theoretical character because of the parts cast in moulds fitted with coolers depending on overall size and shape, provide an homogeneous structure throughout the part thickness with uniform distribution of phases and suitable grain dimension, causing the mechanical properties isotropy.

2. Laboratory experiments

This paper intends to check all these theories in case of several experimental brasses alloyed by Al, Fe, Mn and cast in 90 mm diameter moulds. Experimentally, on A1, A2 and A3 code brasses with dissimilar ratios of α phase (table no. 1), several heat treatment conditions were applied..







37

Table 1

Code Cu [%] Zn [%] Al [%] Fe [%] Mn [%] Σea [%] Phase ratio α[%]

Hardness [HB]

A1 62.34 25.40 4.06 4.17 4.03 12.26 0.0 260 A2 64.30 25.44 3.81 2.93 3.52 10.26 18.6 235 A3 63.17 20.62 4.50 3.86 3.85 12.21 36.0 245

Several heat treatment conditions, type of

structure and recorded values of hardness are given in table no.2.

Heat treatments have been performed in the 17m3 output nitre bar electric furnace from UZINSIDER Engineering Research Laboratory.

Due to an alloying constituent ratio of less than 13% the thermal treated brass structure has

been analyzed according to Cu-Zn constitutional diagram [1, 2, 10], considering that Al, Fe and Mn decrease the solubility of Zn in Cu, dissolve in CVC lattice of β phase and oversaturate it.

Table 2

Brass code Process phase Type of structure Hardness

[HB] ccaasstt iinn 9900 mmmm ddiiaammeetteerr mmoouulldd β’ single-phase 226600

hardening at 800°C, 2 (two) hours; water cooling

β’ single-phase 240

second hardening oversaturated β single-phase 250 third hardening oversaturated β single-phase 270

AA11

hhaarrddeenniinngg aatt 880000°°CC ++ aannnneeaalliinngg aatt 445500°°CC;;

55 ((ffiivvee)) hhoouurrss;; aaiirr ccoooolliinngg

β’ single-phase 225577

ccaasstt iinn 9900 mmmm ddiiaammeetteerr mmoouulldd α+β two-phase’, α precipitates needle-like shape

223355

hardening at 800°C; 2 (two) hours ; water cooling

β single-phase 208

second hardening oversaturated β single-phase 220 third hardening oversaturated β single-phase 240

AA22

hhaarrddeenniinngg aatt 880000°°CC ++ aannnneeaalliinngg aatt 445500°°CC;;


α+β two-phase’, α precipitates needle-like shape

223333

ccaasstt iinn 9900 mmmm ddiiaammeetteerr mmoouulldd α+β two-phase’, α precipitates needle-like and

fine lamellae shape

224455

hardening at 800°C; 2 (two) hours ; water cooling

β single-phase 218

second hardening β single-phase 235 third hardening β single-phase 255

AA33

880000°°CC;; 22 ((ttwwoo)) hhoouurrss ;; water cooling ++ 445500°°CC;;


α+β two-phase’, α precipitates needle-like and

fine lamellae shape

224444

Considering the hammer hardening effects as well as diffusion processes of alloying constituents during heating, it results that heat treatments applied to brasses modify hardness values depending on the

resultant structure. Figure 1 to 8 illustrate the microstructures of thermal treated alloys.



38

Fig. 1 - A1 code. Cast or quenched state β’ single-phase brass.

Ferric chloride attack (×100).

Fig. 2 - A1 code. β’ single-phase brass after the third quenching.


Fig. 3 - A2 code. Cast state α+β’ two-phase brass, α=18.6%,


Fig. 4 - A2 code. Quenched state α+β’ two-phase brass, α=18.6%.


Fig. 5 - A2 code. Quenched and annealed (450°C) two-phase

α+β’brass, α=18.6%, Ferric chloride attack (×100).

Fig. 6 - A3 code. Cast state two-phase α+β’ brass, α=36%.




39

Fig. 7 - A3 code. Quenched state two-phase α+β’, α=36%.


Fig. 8 - A3 code. Quenched and annealed (450°C) state

two-phase α+β’ brass, α=36%, Ferric chloride attack (×100).

AA11 ccooddee bbrraassss –– has after casting, a β’ phase fine grain structure. Subsequent to the first quenching the alloy hardness value decreases by 8%, as a result of growing large grains of β phase (irregular solid solution). Repeated quenching lead to hardness values increased by 12%, supposing to appear the Guinier-Preston areas and Cotrell atmospheres as a result of composition homophase fluctuations [2, 3].

When transients the 450°C temperature (Curie-Kurnakov) the alloy subjected to quenching and annealing slightly increases its hardness due to β’ phase appearance.

AA22 ccooddee bbrraassss –– depicts a hardness value decreased by 11% further to the quenching applied, as a result of putting in solution of α phase.

Repeated quenching has registered hardness values increased by 15% as against the first applied hardening due to the occurrence of atom clusters.

A five hours annealing at 450°C (are of β→β’ regular-irregular transformation) involves hardness value increased by 12% in comparison with the quenched structure, due to a needle-like inter and intragranular α phase precipitation and a growth of β’ regular structure.

The value of hardness obtained in this case is identically as it results subsequently to casting, the alloy showing the same structure of fine grains and needle-like shaped α phase precipitates.

After hardening the AA33 ccooddee bbrraassss has registered a lowering of hardness value by 11% as a result of

putting in solution of α phase. Repeated quenching led to a 17% increased hardness value as segregation areas enriched in alloying constituents occurred (Guinier-Preston areas and Cotrell atmospheres).

A 450°C annealing generates an alloy hardness value equal to the one obtained subsequently by casting. In this case, hardening was achieved both by growing a β’ regular structure as it transients Curie-Kurnakov temperature as well as by integral precipitation of needle-like shaped inter and intra-granular α phase.

Knowing the decisive role of structure upon mechanical characteristics (grain size, phase nature, distribution and ratio) heat treatment results achieved on the experimental brasses have required no determination of mechanical characteristics (these being supposed to be unmodified).

In order to emphasize the heat treatment effects on brass structure morphology, the alloying constituents repartition in basic solid solution as well as occurrence of possible inter-metallic compounds or of Guinier-Preston areas have been analyzed by aid of YXA-5-CM8 type electronic microprobe.

X ray images of alloying constituents composition as well as composition variation configuration have been recorded, both for cast specimens and heat treated ones (fig. 9 to 12).



40

Fig. 9 – Variation configuration of

Fe and Zn elements in oversaturated β single-phase of A1 code brass,

further to the third hardening (×600).

Fig. 10 – Composition image of Fe element in cast state β single-phase of A1 code brass (×1000).

Fig. 11 – Composition image of Fe element in cast

state β single-phase of A1 code brass, further to the third hardening (×1000).

Fig. 12 – Compositions image of Fe element in β single-phase of A1 code brass, subsequently to

quenching and annealing at 450°C (×1000).

Fe Zn

Analyzing the registered images it comes out that distribution of iron in A1 code alloy is modified as follows :

in cast state, it is uniform distributed, showing a tendency for micro-zone clusters;

subsequently to repeated quenching, clusters of this element appear (bright area) as a result of Guinier-Preston area and Cottrell atmospheres or FeZn7 inter-metallic compound existence;

after annealing, the density of these segregation areas is decreasing, without being eliminated.

Further to an analysis of variation configurations of Fe and Zn elements, it is seen that no rising in concentration line amplitude of Zn element within iron segregation areas is registered. This situation proves FeZn7 inter-metallic compound as non-existent.

Under this circumstances, supplementary alloy hardening occurred subsequently to repeated

quenching is caused by the appearance of Guinier-Preston areas and Cottrell atmospheres, confirming, thus, the theory Bauer and Hansen [11, 12] have issued with regard to Fe solubility in α and β’ phases.

They assert that, when two-phase brass cooling occurs the elementary iron rapidly precipitates from β phase, in small crystals, causing grain brittleness.

Further to the experiments performed, it comes out that, these precipitates determine a secondary hardening of β phase, Fe atoms setting up barriers on the way of dislocation movement.

In case of two-phase alloys (A2 and A3 codes) the concentrations line of Zn shows no variations of amplitude in Fe cluster areas, fact that proves FeZn7 inter-metallic compound as being absent in structure (fig. 13). After repeated quenching, the same composition variations for the element were emphasized.

Annealing treatment did not eliminate the grain segregation area.

Considering the reduced solubility of Fe in those two cooling phases, the presence of clusters



41

subsequently to an annealing treatment is explained by low mobility of Fe atoms caused by the presence of dissolved alloying elements in β solid solution (fig. 14 and 15).

Distribution of Al and Mn elements homogeneous for both cast and thermal treated alloys – example illustrated in figures 16 and 17.

Fig. 13 – Variation configuration of Fe and Zn elements in α+β two-phase of A2 code brass,

after repeated quenching (×600).

Fig. 14 – Composition image of Fe element in two-phase brass after the third quenching (×1000).

Fig. 15 – Composition image of Fe element in two-phase brass, after applying quenching and

annealing at 450°C (×1000).

Fig. 16 – Composition image of Mn element in A3 code brass after applying a quenching and an

annealing at 450°C (×1000).

ZnFe

Fig. 17 – Composition image of Al element in A3

code brass after applying hardening and annealing at 450°C (×1000).

3. Conclusions

1. In case of high-tensile foundry brasses it comes out that heat treatments alone cause no significant changes within structure and hardness morphology as against the cast state and mechanical characteristics, respectively.

2. As a result of applying an annealing at 450°C, two-phase brass structure is likewise to the cast state one-fine grains of β’ phase in which inter- and intra-granular α phase is precipitated as needle and fine lamellae-like shape, without any alteration of mechanical characteristics. The hardness obtained in case of two-phase brasses is identically to the one resulted when mould casting.

3. Repeated hardening treatment applied both to single-phase and two-phase brasses generates an



42

additional alloy hardening due to more enriched in Fe segregation areas (Guinier-Preston areas and Cottrell atmospheres).

4. High-tensile brass mould casting procedure is similar to the quenching-annealing treatment. Controlled cooling rate produces a two-phase α+β’ structure with fine grains and homogeneous distribution of α phase, enabling to obtain several higher mechanical characteristics.

References

[1]. Ienciu M., „Smelting and casting of non-ferrous alloys”. The Technical Publishing House, Bucharest, 1982. [2]. Lahtin Y., „Physical metallurgy and heat treatments”. The Metallurgy Publishing House, Moscow, 1984. [3]. Mantea St., “Physical metallurgy”, The Technical Publishing House, Bucharest, 1970.

[4]. Arnaud D., „Traitment par un flux de la crique du laiton”. Fonderie, France, pag. 153, 344/1975. [5]. Whitwham Donald, “Traitments thermiques des alliages de cuivre”. Techniques de l’ingénieur, France, M 1295–10/1989. [6]. Nardin Mario, „Behaviour of special ferrous materials in working conditions”, The Romanian Academy Puiblishing House, Bucharest, 1998. [7]. American Society for Metals (A.S.M.), “Metals Handbook”, USA, vol. 29/1979. [8]. A.S.M., „Source book on cooper and cooper alloys”–sect. VIII. Treatment thermical, USA, 1979. [9]. Kanvenichry I., „A short hand book of heat treatment”, USA, 1987. [10]. * * *, “Properties and Selection of metals”. Metals Handbook, vol. I , Novelty Ohio, USA, 1991. [11]. Titarev N. I., „Mn effect upon Fe alloyed brass properties”. MITOM, S.S.S.R., pag. 30, 3/1990. [12]. Chrzanowski U., Klingelhöffer H., “Structural modifications of non-ferrous alloys”. European Metallographie Conference (EUROMET), Friedrichshafen, Germania, pag. 255, 1995.



OBTAINING AND CHARACTERIZATION OF FERRITE POWDERS

FROM METALLURGYCAL WASTE MATERIALS

Dumitru DIMA

„Dunarea de Jos” University of Galati, Romania E-mail: [email protected]

ABSTRACT

Pickling solutions wasted from iron and steel works contain big quantities of Fe II and Fe III salts. In this work there is considered the superior capitalization of these salts but in the first stage it was started with FeSo4·7H2O got as secondary product in the Pickling process from L.B.R. of S.C. ISPAT SIDEX GALATI company. There is proposed an original procedure of capitalization for this secondary product which can be transformed in ferrite powders having a large practical application: pigment, complementary material in composition, magnetic support, to detect the magnetic fields and screening systems of these, etc. There have been made a compositional, structural and dimensional characterization of these magnetic powders through difractometry Rx, spectroscopy SEM, PFC 200 Particle Counter (ISO 4406, NAS 1638 and ISO 11218).

KEYWORDS: ferrite particles, composite materials, additives composite materials, pigments.

1. Introduction

The superior capitalization of the waste materials from the iron and steel works must be a permanent preoccupation of the research due to the effects over the quality of the environment but also over the dough regarding the decreasing of production costs.

The extended level of production involves some new serial technologies regarding the capitalization of waste materials.

This work has in view these major preoccupations offering valid solutions regarding the capitalization of some solutions reach in FeSO4, resulted from the sheet iron Pickling in the roll process to cold air.

2. Obtaining technology of Fe3O4 from solutions that contain Fe3O4

From solutions of FeSO4 through precipitation with (NH4)2C2O4 we can obtain oxalate of Fe II:

FeSO4 + (NH4)2C2O4 = FeC2O4 + (NH4)2C2O4 It can be thermic decomposed in vacuum at

temperatures over 570º in FeO:

FeC2O4 ⎯⎯⎯ →⎯> Ct o570 FeO + CO + CO2 Fe II oxide is structurally a limit form of iron’s

oxides (FeOx, with x = 1,0 1,5 ) stable at temperatures over 750º.

Under certain conditions of temperature and pressure, the Fe II oxide changes in a mixed oxide Fe II and Fe III SPINEL type:

3FeO + ½O2 = Fe3O4 This oxide represents a limit structure which in

certain conditions can be prevailing. The original technology Fig.1 proposed to get

dominant Fe3O4 has the root in the premise of its thermodynamic stability between temperature limits of 500 - 550º C and of the atmosphere reaction autoreductible consisting in a mixture of CO and CO2.

In order to get a good output in FeC2O4 there must be used a overquantity of 10% of ammonium oxalate, substance that even remains impure in small quantities, do not affect the quality of the final product because under reaction is decomposed thermic as such:

(NH4)2C2O4 ⎯→⎯ Cto

2NH3 + H2O + CO + CO2

43



Fig.1. The technological scheme of Fe3O4 obtaining through the thermic decomposition of the iron

oxalate. Thus, at 550ºC, Fe II oxalate is thermic decomposed in Fe3O4 under oxygen traces:

FeC2O4 + 2O2 (traces) ⎯⎯⎯⎯ →⎯ −= Ct o550500 Fe3O4 + 6CO2

This fact is based on the instability of FeO

under 570º and especially on the presence of the atmospheric oxygen in traces which will finally lead to obtaining of the mixed oxide Fe II and Fe III with good magnetic proprieties.

Fe II and Fe III oxide got through this process is much easier to be obtained presenting a chain of simple and efficient operations.

The granulation of the final product is easier to fix and to control because the iron oxalate can be efficiently ground and the thermic decomposition can loose the final product stopping the formation of crystalline conglomerations.

3. Researches regarding the main

characteristics of the ferrite material Fe3O4

3.1. Elementary characterization

Because the ferrite material has been synthetized having in view a specific technology, there is necessary a characterization of this product regarding the purity and the dimensions of the particles. This process has been realized through difractometry with X rays, through electronic

microscopy SEM and through a special method pointed in cooperation with the chemistry laboratory of “ELECTRICA” S.A. GALATI company.

In the first stage, there has been analyzed the ferrite powder (Fe3O4) obtained through the thermic decomposition of FeC2O4 through difractometry with X rays. (Fig.2)

Fig.2. Difractometric analyze with X rays of ferrite powders (Fe3O4)

The main crystallographic characteristics of the

ferrite powder (Fe3O4)

Dissolving

Dissolving

(NH4)2C2O4

FeSO4 6H2O

Mixture Precipitation Filtration

Washing Drying Grinding Sort out Fe FeC2O4

Grinding Sort out Fe3O4

Ranges of Fe3O4

Thermic decomposition

44



Table 1. shows its high purity level that confirms the theoretical premises from which the

synthesis of this product has been started.

Table 1. The main crystallographic characteristics of the ferrite powder (Fe3O4)

The method could be extended to other types of ferrite, too, depending on the nature of the composite material in whose containing these particles are.

A confirmation in this way is given by the electronic microscopy analyze (SEM) made in the laboratory of CFH L’Ecole Centrale Paris.

In the fig. 3, 4,5,6,7 there are presented images obtained through the SEM technique of the ferrite powders on two types of supports – steel and carbon to various levels of size x3500, x5000, x7500.

Fig.3. SEM analyze of the Fe3O4 powder on steel

support (x7500)

Fig.4. SEM analyze of the Fe3O4 powder on steel

support (x5000)

Fig.5. SEM analyze of the Fe3O4 powder on steel support (x7500)

Fig.6. SEM analyze of the Fe3O4 powder on carbon support (x3500)

N0 Identified stage 2θ [0] d/n [Å] (hkl) Crystallographic system

1. Fe3O4 21,22 4,850 (111) hexagonal 2. Fe3O4 35,00 2,970 (220) hexagonal 3. Fe3O4 41,34 2,530 (311) hexagonal 4. Fe3O4 43,30 2,420 (222) hexagonal 5. Fe3O4 50,30 2,100 (400) hexagonal 6. Fe3O4 62,96 1,710 (422) hexagonal 7. Fe3O4 67,36 1,610 (511) hexagonal 8. Fe3O4 74,04 1,483 (440) hexagonal 9. Fe3O4 84,66 1,326 (620) hexagonal

10. Fe3O4 88,54 1,279 (533) hexagonal

45



Fig.7. SEM analyze of the Fe3O4 powder on carbon support (x5000)

As quality of the image, the sample on a steel

support is better and as concentration of the constitutive elements is normally better the sample on

a carbon support. In Fig. 8. there is presented the SEM Quant analyze with seven characteristic drops on steel support.

The results of analyses on the two types of support (carbon, steel) as well as those realized through difractometry with X rays are displayed in the Table 2.

Analyzing the data from the table 2, there are three limit formulas of the analyzed ferrite for the three cases mentioned above:

I. Fe2,996O4 – through difractometry Rx II. Fe2,64O4 - through SEM Quant on

carbon support III. Fe2,07O4 - through SEM Quant on

steel support

Fig.8. Elementary analyze of the Fe3O4 powder through SEM microscopy on the basis of spectrums of electrons

dispersive energy. (on steel support)

46



Table 2. Elementary analyze of the Fe3O4 powder through difractometric method with X rays and SEM Quant.

SEM Quant Difractometric Rx

Steel support Carbon support N0 Type of the technical element

% massic % atomic % massic % atomic % massic % atomic

1. O 27,59 57,14 32,65 62,86 30,26 62,23

2. Fe 72,41 42,86 67,35 37,14 69,74 39,77

3. Total 100,00 100,00 100,00 100,00 100,00 100,00

It goes without saying the big difference that appears in the case of the formula III in comparison with the ideal limit structure, the influence of the steel support being obvious.

Better results, from a quantitative point of view, can be obtained through SEM Quant on carbon support when the deviation of the formula in comparison with that one ideal limit is of 12%. The error is perfectly admitted for the ferrites used to get composites, having in view that the structure to which is compared is an ideal one. 3.2. Dimensional characterization through

electronic microscopy sem

One could see in the fig. 3,4,5,6,7 that the Fe3O4 powder is as conglomerates from granules of various dimensions but of whose maximum diameter does not exceed 5 microns.

These conglomerates are formed due to the magnetic attraction forces, which through a dry or wet grinding could lead to a optimal granulation for the established reasons of the use.

3.3. Dimensional characterization with

the particles counter (cei 970)

This method is applied to oils for transformer being used to characterize such a product regarding the size of the impure particles.

Particles counting could be realized using an optic microscope and a proper scale of comparison of particles filtered and distributed on a „CAROIAJ” or automatically with an apparatus specialized in counting particles of micronic size.

This apparatus is produces by PALL Industrial Hydraulics Company - New York and is of type PFC 200 Particle Counter. The used working method is in accordance with ISO 4406, NAS 1638 şi ISO 11218 (for particles of 2,5, 10, 15, 25, 50, and 100 microns).

In the first stage there has been realized the counting of Fe3O4 particles using the optical microscope.

For this, it was made a proper dilution of a Fe3O4 sample in a oil for transformer with high purity level.

Before, it was characterized the transformer oil which was of type TR 30.01:

- particles having ∅ bigger than 5 μm = 12.120 -particles having ∅ bigger than 15μm = 847

There have been made two types of dispersion of Fe3O4 in the transformer oil. First type has used a dry grinding of Fe3O4 before to realize the material dispersion.

For this, it was weighted 0,1045 grams of Fe3O4 in the analytic balance and they have been quantitative brought in 800 ml of oil. Due to the big number of particles and to their agglomeration on a „CAROIAJ” filter, there was proceeded to a second dilution, meaning that 2 ml of first solution have been quantitative brought in other 700 ml of oil.

The last one could be “read” to microscope in order to determine the size of Fe3O4 particles, resulting the following experimental data:

- particles having ∅ bigger than 5 μm= 68.824 - particles having ∅ bigger 15 μm = 7.293

From the total number of particles there were substracted the particles accompanying the oil TR 30.01.

Hence, we can draw the conclusion that through a dry grinding one could obtain a particles distribution around 5 μm.

The second type of dispersion of Fe3O4 particles used a “wet” grinding, meaning that it was introduced a small quantity of oil in a mass relation 1 : 1 with Fe3O4 and after that increasing the time with 30 min. there was got a paste from which 0,0206 grams were weighted (0,0103 Fe3O4 dry) and

47



were quantitative brought to the quoted balloon of 100 ml.

0,2 ml were weighted with the dropper from the quoted balloon and it has been brought to the quoted balloon of 100 ml in order to realize a dilution similar to the first trial. (This thing has been made in a different way for the first type to save some oil).

After the reading of the “CAROIAJ” filter to the microscope, there have been obtained the following data:

- particles having ∅ bigger than 2 μm=150.000 -particles having ∅ bigger than 5 μm=43.380 - particles having ∅ bigger than 15 μm =2.351

According to the experimental data, it results that in the case of a dry grinding the ratio between the particles having ∅ within 5 and 15 μm and those ones having ∅ bigger than 15 μm is equal with 9,45 and in the case of a “wet” grinding this ratio rise for about two times, having the value 18,45.

Instead, the outstanding increase in the case of a “wet” grinding is for the particles being within 2 and 5 μm, whose value is for about 150.000.

This fact shows very clear that a “wet” grinding has a remarkable efficiency regarding the dispersion of Fe3O4 particles, in a liquid faze of the experiment realized using the transformer oil TR 30.01.

Also, one could notice that the “wet” grinding lead to the realization of some particles having the size around 1 μm.

This thing has as result the obtaining of some systems Fe3O4 - quasistable oils, meaning that the agglomerations of the particles and their sedimentation can be noticed after 7 - 8 days.

So, the dispersion of a magnetic material such as Fe3O4 in a polymer with a fluidity close to that one of the transformer oils could be easy realized through “wet” grinding.

The grinding efficiency and the dispersion level of the magnetic particles depend on the constructive characteristics of the mill as well as the time of grinding.

There could be realized dispersions of the Fe3O4 particles in laboratory with submicronic particles and a very big stability in time but with small concentrations of Fe3O4 (under 5 % Fe3O4).

For bigger values, the stability of the suspensions decreases very quickly in time, suspensions being stable for a short time, within 5 ÷ 10 hours.

Meantime, the magnetic particles attract each other forming so-called “clusters” that include a fluid in the circular agglomerates, in which they are dispersed [1]

In order to solve such a case, the specialized literature [3] propose the introduction of some

dispersants that avoid the formation of “clusters”, respective the agglomeration and the sedimentation.

Usually, there are use organometallic compounds from the organotetanates category [5].

4. Conclusions

From this work one could notice that from a

solution or an industrial subproduct there can be obtained products of high technology having multiple use at the top industrial fields.

The highlight of this work is the realization of a processing technology of a ferrite material having as input a product considered waste material in metallurgy. The elementary and dimensional characterization of the final product shows its high quality and the efficiency of the proposed method.

Acknowlegments

The author thanks to the laboratory C.F.H. of

L’Ecole Centrale Paris for the SEM analyzes that have been made and to the chemistry laboratory of S.C. ELECTRICA S.A. GALATI for the CEI 970 analyzes of dimensional characterization.

References

[1]Kashevsky B. E. Computer simulation of Ferro suspension structuring in steady and rotating fields, Journal of Magnetism and Magnetic Materials 122 (1993) 34 - 36, North - Holland. [2]Dima D., Doniga A. Studii referitoare la obţinerea uni material compozit fibră de sticlă - poliester - Fe3O4 cu proprietăţi magnetice, Sesiunea ştiinţifică de metalurgie - în volum Galaţi Octombrie 2000 p. 58 – 63. [3]Andrei C., Andrei G. Performanţe în poliolefine, Editura Zecasin, Bucureşti, 1998, p. 270 - 276. [4]Doniga A., Dima D. Rapid investigation method of silicon steel magnetically domains, Proceedings of the 39th International Seminar on Modeling and Optimization of Composites - MOC’39 Sensible experiment in materials Science 26 - 27 April 2000 Odessa. [5]Katz H. S. Haudbook of fillers and reinforcements for plastics, vol. 4, editura Van Nostrand Reinvold, 1978. [6]Luca E. Ferofluide şi aplicaţiile lor în industrie, Editura Tehnică, Bucureşti, 1978. [7]Mc Caig M. Permanent Magnets in Theory and Practice, John Wiley and Sons, Inc., New York, 1977. [8]Muşat V. D. Teză de doctorat - Cinetică heterogenă în sisteme generatoare de oxizi micşti cu proprietăţi magnetice, Universitatea din Bucureşti, 1996. [9]Parker R. J. Advances in Permanent Magnetism, John Wiley and Sons, Inc., New York, 1990. [10]Pelletier S., Gélinas C., Chagnon F. Effect of compaction temperature on magnetic properties of iron - resin composites, Industrial Materials Institute - Québec Canada J3R4R4 - 2000. [11]Pont R. P., Krishna R. M., Nagi P.S. XRD, SEM, EPR and microwave investigations of ferrofluid - PVA composite fillers, Elsevier - Journal of Magnetism and Magnetic Materials 149 (1995), 10 - 14.

48



49

THE PLASTIC DEFORMATION OF STEEL SHEETS FOR

ENAMELING AND DEFECTS OF ENAMELED LAYER

Petrică ALEXANDRU

“Dunărea de Jos” University of Galaţi e-mail: [email protected]

ABSTRACT

The paperwork presents at first in principle enamelling technology for

products obtained from sheet-iron then is described the mechanism formation of defects called " fishscales", the manner of act avoided the appearance, method of measure the sensibility for this defect and the own researches: the connection among the degree of cold-working and the “fishscales” probability appearance.

KEYWORDS: steel enamelling, fishscales, deep drawing steel

1. Introduction

Enamelling is a hot coating process (800-8500C) which consists in melting an enamel powder deposited on a sheet, in such way to obtain a vitrified coating on the steel surface. The obtainable composite material is characterized by a high abrasion resistance, a high chemical inertia and an excellent heating resistance. With these properties, the enamelled steel sheets are used in numerous fields: household appliances, cook ware, sanitary appliances, architecture, chemical industry, etc.

The metallurgists for the role mainly know the hydrogen in the metal brittleness, but it can lead to other drawbacks. During the specially enamelling of extramild steel sheets, it can contribute to surface defects called “fishscales”. For enamellers, it is the most dangerous defect because it can appear after a long delay when the parts are in use. To avoid such a problem, is better to use special steels that are tested for the sensibility among “fishscale” defect. In this article it shows the experimental determinations of influences which the degree of deformation applies to sheet metals, obtained by cold rolling and enamelled, has it about frequencies of "fishscales defect appearance".

1.1. Enamels

The enamels are glasses aboundant in flux elements, with a low transition temperature. Thuse are obtained by melting mineral raw materials and continuously solidified like steel before being broken in fragments.

Two kinds of enamel exist: -the “ground coat enamel”, dark in color, which

ensure adhesion steel-enamel, by presence of oxides NiO, CoO, etc.

-the “cover enamels” applied with or without “ground coat enamel” and which give the protection and decoration.

The table 1 shows two examples by chemical compositions for enamels:

Table 1. Examples of enamels composition

Component Ground [% weight]

Cover (white)

[% weight] SiO2 50 45 TiO2 - 18 ZrO2 2 - P2O5 1 2

BB2O3 14 15 Al2O3 3 2 CaO 4 - BaO 5 - Na2O 11 10 K2O 2 6 Li2O 2 - NiO 1 - CoO 0,5 - CuO 0,5 -

F 2 2 Total 100 100

The preparation is done by casting slip through

the addition fictile, electrolytes and water during the triturating powdery or by dry triturating. The enamel is deposited on support through like modes used in the painting: immersion, simple pulverization, pulverization in electrostatic field, electrophoresis, etc. Before in furnace vitrification, the layers deposited by liquid enamels must be advanced dried.




50

1.2. Proceedings of enamelling

Classical enameling This method demand a good and complete

preparation of support surface: degreasing-washing, pickling-washing, neutralizing and dried. The next step is applied and burning ground-enamel which besides adhesion enamel-steel, have a function in combating the bad effect of reaction gases. For special requirements of aspect, it is proceeding to apply and burning a next layer, by type cover coating.

Directly enameling The layer of ground-coat enamel is

eliminated and on the support is applied only the cover-coating enamel. More economical, this proceeding required instead a steel-sheet decarburized (C free<0,004%), a more advanced pickling and a chemical nickel plating.

Enameling in two layers-an only burning The proceeding was adjusted recently, and

assumed the deposit on the support, whence is in previous worked just through a simple picking-washing-dry operation, of a fine layer by special ground-coat enamel, there to is deposited and the layer of cover-coating enamel, without intermediate burning.

1.3. Formation of “fishscales”

The physical and chemical phenomena implied in vitrifying of enamel layers and the role of hydrogen in these processes are presented in following rows.

Diffusivity of hydrogen in steel From the table 2 it observe the very different

mode of hydrogen diffusivity dependence on temperature in comparison with others elements whence equally dissolved interstitial (C and N) in pure iron. The hydrogen diffusivity has a reduced variation on temperature and remained with a high level for the room temperature.

Table 2 .Diffusivity of hydrogen, carbon and asset as a function of temperature

Diffusivity [cm2/s] Temperature [°C] H C N

20 1,5x10-5 2,0x10-17 8.8x10-17

200 0,55x10-4 - - 700 4,9,10-4 6,1x10-7 4,4x10-7

774 2,75x10-4 - -

Solubility of hydrogen in steel Figure 1 shows the rapidly increasing of the

atomic hydrogen solubility into steel in function of temperature, from very low value at room temperature. It is remarkable the important variation of solubility at allotropic state changing of pure iron.

FeαFeδFeγ

Temperature [0C]

0 600 900 1200 1500 1800300

Solu

bilit

y of

hyd

roge

n [m

l H2/1

00g

Fe]

5

35

30

25

20

15

10

Fig. 1. Solubility of hydrogen in pure iron fortemperature variation.

00

900

800

700

600

500

400

300

200

100

1 2 3 4 5 6 7

Time [min]

Tem

pera

ture

[0 C]

Melting enamel

Impermeable enamel (reactions oxides-enamel)

Porous enamel (steel oxidation)

Vitrifiedenamel

Stoppedreactions

Fig. 2. Thermic cycle for enamelling productsfabricated from steel sheets.



51

Thermal cycle of enameling The heating and especially cooling according

with figure 2 are very rapidly. Mechanism of defect formation It can differentiate three stages of fishscale

formation. The first two stages correspond with heating and respectively cooling and the last stage correspond at room temperature processes.

The heating is usually made into atmosphere abundant in water vapors which partially has one origin at water evaporated from enameling layer (deposited from slurry immersion) and it produce the oxidation of steel support. The oxidation produces atomic hydrogen, which penetrate very easy in interstices of the steel crystal lattice by reason of high solubility level.

The relatively speed cooling produce the solidification and formation of vitrified enameling layer which is impermeable for the hydrogen and such as to this gas remain “prisoner” in the steel.

To room temperature the hydrogen diffusivity rest at high level and this is a favorable factor for hydrogen accumulation by molecular form at all the steel imperfections, especially to the interface enameling-steel. The local pressure can have very high values and leads to break the enamel with formation of the semicircular defects. This defects is called “fishscales”.

1.4. Cures for “fishscales”

The high hydrogen affinity of the steel to high temperature and even the impermeability of enamel for this gas to room temperature are the principal factors of fishscales formation. For avoid formation of these defects it can apply two type of principal measures.

Non-metallurgical measures Measures refer to restrictions of atomic

hydrogen produced in course of enameling cycling firing:

-utilization the enamels with low content of OH groups;

-effective and advanced drying after enamel applications from slurry or other forms of enamel watery applications;

-maintenance a low level for the dew-point temperature in furnace atmosphere.

Equally, most be operated for the optimization of fabrication process such as to obtain a very good adherence enamel-steel, and a great mechanical resistance for the vitrified enamel layer.

Metallurgical measures Consists in increase of hydrogen absorption

capacity and maintenance the gas in steel. For this it make so-called “hydrogen traps” in the metallic material.

Three types of the “traps” which could catch the hydrogen are occurred in steels:

-cavities near the interface metallic matrix and the inclusions or precipitates;

Theses are make by means of cold-rolling that could lead to formation nearly the very hard inclusions (Al2O3) but even for the soft inclusions (MnS) to a large number of micro-cavities capable to accumulated the molecular hydrogen.

-cavities that are make by broken the aggregates of cementite particles;

A high level of coiling temperature after hot rolling can lead to make the “islands of cementite” formations which by next cold-rolling operation are broken and generate a great number of micro-cavities.

-“chemical traps” Theses traps are occurred only in case of allied

steels, which at high temperature, during the enamel firing, can be generate the stable chemical compounds (hydrides).

1.5. The sensibility test for “fishscale” Many methods have obtained for attempts and

selects the steels in report with the out of order formation of enamels. Among these, the standard EN 10209 elaborated relatively recent, recommends two methods of characterizes the susceptibility to the formation "fishscales", corresponsive for two types of tests most used-up.

Method by enameling of steel samples In this case, the conditions of testing is most

severe in report with industrial practice in such kind that eventually defects let us appears quick, thus:

-use a special enamel, appoint" sensitive to fishscales";

-the dew of atmospheres from furnace is to a high level;

-the enamelled samples are accelerated aging; Directly steel testing Is measured directly the catcher capacity of

hydrogen in steel. This method has the advantage that estimates the resistance to the formation fishscales without else takes count of the whence incident can step in progress operations of enamel. The principle consists in the production, through electrolysis, hydrogen on a sample face of sheet metal and the measuring on the other face amount of gas whence cross the sample, depending on time.

2. Experimental conditions

The sheet metal by rimmed steel for deep drawing type A3, (STAS 9485-80) produced to SIDEX S.A. Galatzi, with thickness of 1mm was used from samples. By cold-rolling with degree of



52

decrease: 5; 10; 20; 30; and 40% was obtained respective samples with the thickness: 0,95; 0,90; 0, 80; 0,70; and 0,60 mm;

The cycle of processing whereat were submissive samples is presented in figure 3. The firing of samples for the enamel formation layers are made in electric furnace.

After the accelerated aging, were counted the fishscales on the surface samples.

Samples80x80x1mm

Cold-rolling

Ungreasing

Drying in stove

Samples immersion in enamel slurry

Firing in furnace8500C, 5min

Aging in stove800C 48h

Drying in stove

Fig. 3. The block diagram for samples enameling.

3. Results

The variation “fishscales” number depending on the degree of deformation through cold rolling is presented in the table 3, and the graphic representation in figure 4.

Micrographies from figure 5 present the enamel structure, whence isn't absolutely compact, only that the spherical cavities insulated by-path and ergo the enamel is impermeable, and the test is relevant.

0

2

4

6

8

10

12

14

0 10 20 30 40 5

Deformation degree [%]

0

Fish

scal

es\d

m2

Fig. 4. The deffects type fishscales depending of coldrolling degree.

Table 3. The “fishscales” number depending on

deformation degree.

Deformation degree

[%] Fishscales/dm2

40 11,5 30 13,0 20 4,8 10 0 5 0 0 0

Fig. 5. Aspect of enamel vitrified layer, (x300).

4. Conclusions

General conclusions 1. To enameling products from steel, the

defects fishscales appearance is determinate of the hydrogen presence in material metallic structure.

2. The presence of the water vapors in atmosphere of enameling furnaces permits the



53

formation of atomic hydrogen. Due to very little sizes, the atomic hydrogen is able to have dissolved in considerable amounts in steel, especially after the steel transformation in austenite.

3. For avoided the appearance defects "fishscales", can action in main on two ways:

-limitation the volume of hydrogen from atmosphere furnace during enameling process (burnt gas) and from used water for liquid of enamel slurry, by immersion of products;

-absorbedness steels for hydrogen; 4. The European standard EN 10209

recommends two methods of testing for sensibility steels against the defects appearance" fishscales", the first requires enamelling and the second allow the mensuration directly capacity steels of hydrogen inclusions in the structure.

Conclusions obtained from the experiments 1. Cold-working has a major influence about

defects appearances" on product from enamelling steels;

2. In the case of band from steel mark A3, for which accomplished the researches, it consisted that

the domain of deformation degrees 20-40% must avoided, in practice of deep-drawing produced what were submissive enamelling, because the susceptibility of appearance the defects type fishscales is very high;

3. The testing practice for sheets and bands of steels in delivery state must be improved;

4. It’s recommended, for the sheet metals and bands testing for sensibility to fishscales, let us is done after plastic deformation previous, in a large range of deformation degrees and by tensional and deformation of the scheme types, that could appears in the technological process to beneficiary.

Refrences [1]. P. Hemmen, V. Cholet, L. Guillot,, L’hydrogene et les aciers pour emaillage, La Revue de Metalurgie-CIT, nr.1, 1995; [2]. G. M. Pressouyre, Trap theory of hydrogen embrittlement, Acta Metalurgica, vol. 28 (1980), p. 895-911. [3]. The EN 10209 standard.



54

SURFACE HARDENING BY MEANS OF MICROALLOYING

AND VIBRATORY ELECTRODE DEPOSIT

Simona BOICIUC


ABSTRACT

Alloying and deposit of metals and alloys by means of electric spark with

vibratory electrode is a efficient new effective method to change chemical composition, structure, wear and corrosion behavior of steel material surface layer.

KEYWORDS: vibratory electrode, wear resistance, surface hardening

1. Introduction

The life span of a machine, plant or equipment depends on the rate their component parts are wearing. As a matter of fact, wear, fatigue and corrosion phenomena are the most aggressive factors leading to placing out of service of a part.

Surface engineering, as a technical interlinked study field science is a new later years concept which came out in progress industry countries as a result of spectacular development of surface treatment.

Tom Bell from Birmingham University, England, has given a definition to this new concept: “surface engineering involves, in principle, design of surface and sublayer taken together, as a system, in order to grant performances neither surface nor underlayer is capable separately”.

Therefore, surface engineering is not only a simple choice for one of the existing surface treatment technologies, but it refers to design of base material-surface layer system so that it should give a very well response to its functional role on an efficient use of materials at high accessible price costs.

Thus, electric spark alloying and deposit method is coming to complete the classical thermochemical and thermal processing procedures as well as the unconventional thermal processing procedures contributing to an improvement of part surface layer properties.

2. Experimentally The principle of metal part hardening by means

of electric sparks consists in the fallowing : in case of air spark discharge under the action of rectified pulsatory current, a polar transfer of electrode (anode) on part surface (cathode) takes place. This material alloys the part layer and, in a chemical association with air dissociated atomic nitrogen, carbon and part material make a hardened diffusion layer and a high wear resistant white deposited layer, both shoring complex chemical combinations of nitrides, carbonitrides, high stable carbides as well as hardening structures.

Figure no. 1 shows a principle circuit diagram for electric spark alloying (deposition) installation.

R+

-

1

23

4

5

Fig. 1. Principle circuit diagram of the electric spark alloying and coating system

1 – power supply unit; 2 – vibrating device; 3 – electrode (anode); 4 – part (cathode) 5 – contact plate.



The main requirement of this method is to

determine the electric conductivity for part and electrode, as well.

The most important characteristic of this procedure is mainly the energy necessary to discharge electric spark within the interelectrode gap. Therefore, when electrode approaches near the part, at a disruptive interval of about 10 μm, continuous electric pulse discharge take place and it ends when electrode contact occurs.

After disruption of electrode gap with a condenser stored energy, surface leading to an electric corrosion of there, an electrode material polar transfer of electrode on part surface develops and a surface layer with suitable mechanical properties is built-up.

Discharge energy varies from 8 to 18 Joule at supply voltages within 15 to 220 V; the average current enables control from 0,2 to 80 A depending on the characteristics of alloyed and deposited layer and of gained surface quality.

Also, it have been found that the smaller discharge energy can be the smaller deposited layer and its gained surface roughness are; the deposited layer being more dense and with a cleaner surface. In general, thin deposited layers are preferable to those thick and porous and usually crack ones.

Surface part layer processes developed when deposition and alloying with electric sparks given by a vibratory electrode occurred are the following:

• High speed (ultra rapid) hardening of steel and cast iron liquid and solid state due to ultra rapid cooling of heated surface, the heat being taking over by the wide mass of cold cathode metal. During short (10-3 to 10-5 sec) electric discharge, high temperatures of 5000 to 11000 ˚C/sec are developed which melt and local evaporate the material located in the surface layers of part and electrode, as well. High speed cooling (104 to 105 ˚C/sec) results in a specific hardening structure which imprints special surface layer properties.

• Steel carburation carried out by adsorption and diffusion of active carbon electrode atoms or from interelectrode medium into the base layer when very

fine and uniformly spreaded carbides and carbonitrides are built-up.

• Nitriding is achieved by absorption, dissociation and diffusion of air molecularly nitrogen during electric discharge when fine iron, chrome, aluminum, etc. nitrides are formed, playing a special role of hardening and non corrosive protection.

• Alloying and deposition – electrode alloying constituents, by contact or dripping off under polar action, are absorbed and they diffuse over the melted surface and in the solid layer, later on providing an area of deposition and alloying material.

Alloying involves increased values for resilience, corrosion resistance, steel hardening capacity properties and alteration of electric and magnetic areas.

Material and electrode section are chosen dependent on material, part type and purpose to be fulfilled.

Characteristics of layer deposed by alloying and deposition with electric spark and vibratory electrode method depend on the chosen mode of operation, nature and section of electrode, nature of part material and purpose to be achieved.

This paper describes the process of hardening for several grades of steel: OLC 35, OLC 55, 15Cr08, 30Cr130 with distinct electrodes:

- metal sintered carbides; - native metal: W. Metal sintered carbides electrodes are usually

used to increase durability and hardness showing the following chemical composition: Ti15Co6 – WC-79%; TiC-15%; Co-6% and WCo8 – WC-92%; Co-8%. Wolfram electrodes are used especially to get increase value of durability and refractoriness.

Chemical composition of the aforesaid steel grades was spectrographically determined by means of a DV6 type of spectrometer and is given in table no. 1. The value of the pulse current is chosen depending on the nature and properties of alloyed and deposed layer, while the amplitude of electrode vibrations is established at its minimum, conditioned by prevention of electrode – part sticking.

Table 1 Chemical composition of examined steel

Chemical composition,% Steel grade C Si Mn S P Cr Ni Al

OLC35 0.34 0.31 0.68 0.016 0.012 0.13 0.186 0.030 OLC55 0.54 0.21 0.75 0.030 0.019 0.185 - - 15Cr08 0.15 0.23 0.96 0.013 0.016 0.79 0.10 0.029

30Cr130 0.281 0.275 0.299 0.001 0.023 12.87 0.30 0.010 Alloying and electric spark deposition were

carried out with the aid of Elitron type plants – a product made in Moldova Republic by Experimental

Works of Science Academy, available on

Uzinsider Engineering Galaţi site.

55



56

Table no.2 depicts relevant running the samples

were subjected to as well as the related values of hardness (HV5 – 5 kg load) achieved further to depositions.

Table 2 Treatment conditions applied to steel specimens

Operating parameters Steel grade Electrode Running Amplitude

Hardness of base material HV5[daN/mm2]

Hardness of treated material HV5 [daN/mm2]

Ti15Co6 3 8 180 233 WCo8 5 8 180 240 OLC35

W 3 5 180 253 Ti15Co6 3 8 224 245 WCo8 3 8 224 253 OLC55

W 3 10 224 274 Ti15Co6 3 8 168 210 WCo8 3 8 168 224 15Cr08

W 3 10 168 233 Ti15Co6 3 8 249 260 WCo8 3 8 249 315

30Cr130

W 3 10 249 317

0

50

100

150

200

250

300

350

Ti15

Co6/

OLC

35

Ti15

Co6/

OLC

55

Ti15

Co6/

15Cr

08

Ti15

Co6/

30Cr

130

WCo

8/O

LC35

WCo

8/O

LC55

WCo

8/15

Cr08

WCo

8/30

Cr13

0

W/O

LC35

W/O

LC55

W/1

5Cr0

8

W/3

0Cr1

30

Har

dnes

s [H

V]

Fig. 2. The hardness obtained as a result of microalloying

Further to a thorough analyze of the results

obtained in table no.2 and figure no. 2, it derives that the applied running fulfill the task of hardening the experiment used steel grades. WCo8 and W electrodes involve higher values of hardness, thus, for the same basic material, the deposed layer hardness is growing in the same time electrode hardness grows.

One may also observe that treated steel hardness is greatly influenced by the chemical basic composition. Therefore, the more numbered and in a larger amount the alloying constituents are – coupled with a greater preliminary hardness, the higher deposed layer hardness should be. This being caused by chemical compounds, nitrides and carbides built-up in layers.

Metallographic analyses applied on hardened specimens shown deposed white layers, clearly or direly boundered depending on the treatment applied

version and nature of electrode involved. Deposed layer microstructures obtained by means of various treatment versions are shown in figures no. 3, 4, 5 and 6.

Upon a close examination of figure no. 3 (OLC 35), it observes that depositions with W and WCo8 electrodes present relative homogeneous, continuous, well stiffed and adherent white layers.

Further to a thorough analysis of figure no. 4 (OLC 55), Ti15Co6 electrode deposition provides a white, homogeneous, continuous and well stiffed layer, showing the maximum thickness.

Ti15Co6 and WCo8 electrodes deposits involved for 15Cr08 steel grade present white, relative homogeneous, adherent and continuous layers while their thickness values are smaller than in case of W electrode deposit showing a higher average thickness and a diminished stiffening (figure no. 5).



57

Ti15Co6 electrode deposit related to 30Cr130 steel grade shows a white, relative homogeneous and compact layer unlike WCo8 and W electrode depositions, less homogeneous and compact layer

unlike WCo8 and W electrode depositions, less homogeneous and compact, respectively (figure no. 6)

Ti15Co6 deposit

Average layer thickness 58,6 μm Magnify x500

WCo8 deposit


W deposit


Fig.3. Microstructures of deposed layers on OLC 35 steel , nital 2% etched



58

Ti15Co6 deposit

Average layer thickness 72.5 μm Magnify x500

WCo8 deposit

Average layer thickness 46μm Magnify x500

W deposit

Average layer thickness 63 μm Magnify x500

Base material Magnify x100

Fig.4. Microstructures of deposed layers on OLC 55 steel, nital 2% etched



59

Ti15Co6 deposit


WCo8 deposit


W deposit



Fig.5. Microstructures of deposed layers on 15Cr08 steel, nital 2% etched



60

Ti15Co6 deposit


WCo8 deposit


W deposit



Fig.6. Microstructures of deposed layers on 30Cr130 steel, electrolytic proceeding 50% HNO3



3. Conclusions

By comparing with other methods, the alloying

and deposition by electric spark presents a series of advantages: the deposed metallic layer presents a resistant connection with the basis material; the method makes possible the deposition of pure metals (Ni, Cr, Mo, W, Ti) or metallic alloys; it is not necessary a preliminary preparation of the deposition surface.

Ti15Co6, WCo8 and W electrode deposits have caused an increase in hardness of steel grades involved in this experiment, due to the built-up surface layer composed of carbides and complex nitrides.

Treated steel hardness depends in a large measure on chemical composition of basic material. Thus, the more and in a larger amount the basic material alloy constituents are and with a higher value for its preliminary hardness, the higher value of deposed layer hardness will be.

The quality of hardened layer also depends on electrode material composition, so that, WCo8 and Ti15Co6 alloys are recommended for member hardening, while Ti15Co6 is recommended for electrodes used for hardening of cutting tool edges.

A real advantage for steel processing with electric sparks consists in the fact that surface processed in this way develops an improved hardness, durability, heat as well as corrosion resistance.

Metallographic analyses performed on hardened specimens have outlined hite deposed layers, clearly or direly bordered, depending on treatment applied version and electrode nature involved in this experiment.

Vibratory electrode alloying-and-deposit method may be applied on hardening of synthetic fiber knives, strip cutting cylinders and rollers, cross-cut pipe disk cutters (Ti15Co6 and Wco8 are used as electrodes). Further to the application of this deposits, the aforesaid tools grove increased values of endurance, durability and hardness, this providing a favorable influence upon efficiency.

References

[1]. A.T. Busik a.o. – Versions of surface alloying method using electric sparks, Elektronaia obrabotka materialov, nr.1/1991, pag. 22-25. [2]. V.I. Ivanov – Material hardening using electro-erosion alloying, Kuznecino stampovo-cinoe proizvodstvo, nr.9/1990, pag. 11-12. [3]. Contract 94C/B12 – Researches concerning the improvement of corrosion and ear behaviour of iron and steel metallurgy parts using electric spark deposit procedure, ICPPAM Galaţi, 1994. [4]. Billy J. Zabavnikv – Treated tools steel investigation with electronic microscope, Koveve materialy nr. 2/1985, page 166-173, Cehia [5]. Adrian Alexandru, a.o. – Metallic carbides layers with high abrasive wear resistance deposit on alloyed steels with duplex treatments, National Conference of Advanced Technologies and Materials, page 42-46, Galaţi, Romania.

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62

MATHEMATICAL MODEL FOR THE SINTERING

IRON-GRAPHITE POWDERS MIXTURE SIMULATION

Ionel PETREA

“Dunărea de Jos” University of Galati e-mail: [email protected]

ABSTRACT

The simulation of the sintering iron-graphite powders mixture allows to

establish soaking time for obtaining sintered pieces with maxim resistance to wear, of 0,7-0,8 % carbon percentage adequate to the eutectoide composition.

It is presented a computational integration model of diffusion equations, using the finite differences method.

KEYWORDS: pearlitic structure, free cementite, soaking time at sintering,

antifriction material

1. Introduction

The sintering of the iron-graphite system is characterised by the diffusion of a part of graphite, which forms a skeleton with a pearlitic structure. In figure 1 it is presented a model for the graphite dissolution into the iron metallic mass.

The iron powder carbiding can be realise by two mechanisms (fig.1):

a) The aerifying theory - The graphite aerifying C + 1/2 O2 = CO - The CO adsorption on the iron surface - The cat-cracking 2CO = CO2 + C - C diffusion in Fe cell - The CO regeneration using the reaction: - CO2 + C = 2CO

b) The diffusion theory- the direct C inlet in the Fe cell by diffusion in volume

Fig.2 The free standard enthalpies temperature variations of the C – O system reactions

In figure 2 are presented the free standard

enthalpies temperature variations of the C – O system reactions . The ΔG0 values depending on T for the Belle-Boudouard reaction are represented by the line 4. At temperatures higher than 973 K, the free standard enthalpies temperature variations reaction is positive. Consequently, in the presence of a sufficent carbon quantity, over these temperatures, the reaction takes place, forming CO. In figure 3 are presented the equilibrium conditions of the Belle-Boudouard reaction depending on temperature. At temperatures

2

1

4

3

+100 +50 0 -50 -100 -150

ΔG0

,kca

l

1000 1500 2000 2500 Temperature, K

Reaction 1 2 C + O2 = 2 CO Reaction 2 C + O2 = CO2Reaction 3 2 CO + O2 = 2 CO2Reaction 4 2 CO = CO2 + C

CO

Graphite

Fig.1 Model for the graphite dissolution into iron

Iron grains



63

higher than 1100 oC, in the presence of the solid carbon , the CO2 is missing (the whole CO2 quantity is transformed in CO).

Contrary, at low temperartures, CO, not being a

stable combination, is decomposed (below 400 oC in the system exist only gaseous CO2 and solid C). Consequently, at high temperature, CO has a higher thermodynamic stability, whereas at low temperatures, CO2 is more stable than CO.

Considering this termodynamic aspects, results that the carbon dissolution into iron during the sintering iron-graphite mixture powder is realised at low temperatures through the decomposition CO in CO2 and carbon, and at high temperatures, through the diffusion in volume of the carbon in iron particles.

These chemical reactions are not easy to control, so it can lead to variable results from one burden to another, especially in the case of the products that contain more than 3% graphite. As a result, the pearlitic-graphitic formation is very difficult to be made, because at the sintering temperature of the components at 1050…1100 °C, the austenite is capable at dissolving more C than of making a pearlitic structure.

To avoid the secondary cementite forming in the structure had been proposed other technological alternatives, for example the mixing of iron and graphite powder with copper powder.

At 1050°C, which represent the usual sintering temperature of the iron/graphite mixtures, the quantity of dissolved carbon is 1,7%. When an austenite with this composition is decomposed during the cooling process, the resulted structure consists of pearlite and about 15% cementite. To assure a graphit-pearlitic structure, it is necessary, independent of sintering temperature and the quantity of graphite from the mixture, to maintain the dissolved carbon at about 0,8%.

2. Mathematics model

The mathematics model considers that the iron-graphite powders mixture sintering is a sintering in

solid step with a gas-solid reaction similar carburizing.. It is supposed that the graphite is relatively uniformly distributed and plays a role of case-hardening compound.

100 75 50 25 0

The purpose of the simulation is the determination of the sintering temperature and of the soaking time for obtaining sintered pieces with maxim resistance to wear, of 0,7-0,8 % carbon percentage adequate to the eutectoide composition, knowing the physics-chemical characteristics of iron and graphite powder.

Also, it can be considered that the mass transfer is realised in one way and the powder form doesn’t affect essential the process evolution. (is considered only the X axis.).

It is presented a computational integration model of diffusion equations, using the finite differences method. The carburized layer profile is determined using iterative calculus of differences, through the integration of the second Fick equation.

The boundary condition for the interface case-hardening compound-iron powder, is determined using first Fick equation, integrated with Runge-Kutta method order IV.

It will be considered the temperature influence and of carbon concentration over the diffusion coefficient.

Mathematically, the second Fick law is expressed

using the differential equation order II: with initial condition: C(0,x)=C0 , where x∈[0, L/2]

and the boundary conditions : C(t,L/2)=Ci

C(t,0)=CS , for t∈[ 0,tmax] in which CS is the solution of the differential equation order I (the first Fick law):

with initial condition: C(0)=C0

Notations: C(t,x) = the carbon concentration realised after time t and x depth, [%]; t = the concentration evaluation time[min]; x = the concentration evaluation depth with values between 0 and L/2; L = the iron powder grain dimension [mm]; C0 = the initial carbon concentration in the iron powder[%];

2CO→CO2+C

0 25 50 75 100

600 800 1000 Temperature, oC

% C

O % C

O2

Range I

Range II

CO2+C→2CO

Fig.3 Equilibrium conditions of the Belle- Boudouard reaction

(1) ( )

xCD

xtC

C ∂∂

⋅∂∂

=∂∂

dxdCDc)CCe(K =− (2)



64

Ci = the carbon concentration imposed after sintering Ce = the equilibrium carbon concentration depending on the graphite; CS = the concentration after time t at the medium-grains interface; K = the adsorption speed carbon constant at the grains surface : K=K0 exp(-Qa/RT) , [cm/sec] where: Qa = the activation energy of the adsorption process; R = the perfect gas constant ; For iron : K=0.00136 exp(-5570/T), [cm/sec]; Dc = the difussion coefficient Dc = D0 exp ( -Qd / RT) Qd - the activation energy of the C difussion into austenite (123.920 J/atom gram) D0 = a+b %C, where a=0.07, b=0.06;

The term from the equation (1) can be written: Because :

the equation (1) becomes:

The equation solution with the finite differences method (the networks method) suppose the determination of the situation correspondent xj position at time ti+1 using three points situated at the distances xj-1, xj, xj+1 at the time ti . The network step for the distance hx is fixed, its value influencing the calcul precision. The time step ht is variable, depending on the effective carbon procentage concentration. The realisation of the computational stability (minimal integration error) impose the condition for the ratio between the two steps :

For the ratio σ is recommended the value 1/6 which corresponds to the minimal integration error. Considering the diagram with finite differences and noting ,

it will result :

As a consequence of the C concentration permanent using in the gas – metal interface, the diffusion coefficient will rise and the integration time step will decrease continually. After every determination concentrations step at ti time, using the relation (4), for x=0, the C concentration value C(t,0) = CS ., in which CS is the differential equation solution order I. Using this value is being calculated Dc and than the new step value h t.

c2xt D/hh σ=

( )CCDk

dxdC

e −=

Differential equation order I: can be integrated with the algorithm Runge –Kutta

by order IV:

( )6

kkk2kCC 4321

1,i0,i+++

+=

where: k1=ht F(Ci,1)

k2=ht F(Ci,1+k1/2) k3=ht F(Ci,1+k2/2) k4=ht F(Ci,1+k3) where function F has the form F(u) =hx k(Ce –u)/Dc.

3. Calculation algorithm 1. Are introduced the initial dates : the sintering temperature T, the powder mixture C concentration (Ce), the initial C powder concentration (C0), the activation energies of the adsorption and diffusion process (Qa, Qd), the perfect gas constant (R), the constants k0,a,b, the raport value σ, the diameter grains L, the number of points, in which the calcul will be made with finite differences n, the C concentration impose after the sintering Ci.

2. Is initialised: - the time at the value 0 - the calcul step on direction x: hx=L/2 n - the time stept for the first iteration: - the matrix in which is stocked the concentrations value with the dimensions 2*n+1, with the value C0 ;

3. It is calculated at the every iteration : - the momentary value of the diffusion coefficient: - the momentary value of the time step: - the C concentration values 1..n-1, with the relation (4) -(the values will be stocked in the second line of the concentration matrix) ; - the concentration value at x=0, with the relation (5) ; - the new value of the time step; -the new value of the time equal to the previous value plus the new time step;

4. It is verified if it has been reached the C concentration impose after the sintering ; if it hasn’t been outrun, is repeated step three, and if it has been outruned , it is shown the sintering time.

)RT/Qdexp(bC

Dc−=

∂∂

2

22RTQd

xCDc

xCbe

tC

∂

∂+⎟

⎠⎞

⎜⎝⎛∂∂

=∂∂ −

c

2x

t2x

ct

Dh

h21

hDh

σ σ=⇒<⋅

=

( )jij,i x,tCC =

(3)

2

22

xCDc

xC

CDc

xCDc

x ∂∂

+⎟⎠⎞

⎜⎝⎛∂∂

∂∂

=⎟⎠⎞

⎜⎝⎛

∂∂

∂∂

( )( )]CC2C

CCbCab

41[CC

1j,ij,i1j,i

21j,i1j,i

j,ij,ij,1i

−+

−++

+−

+−+

σ+=

(4)

(5)



65

The source programme written in C++ using the previous algorithm, calculates the soaking time at sintering , depending on temperature , the powder granulation, the graphite mixture proportion and on the C percentage dissolved in the iron mass, the last value being imposed by the user.

The calculations are iterative realised, at every iteration are presented the diffusion coefficient value, the time step value and the C percentage distribution dissolved in the iron particles, from the iron-graphite interface up to the granule centre. An example of values obtained after the programme rolling is shown in table 1.

Table 1 Values obtained after the programme rolling

Temperature at sintering , oC

Powder granulation, μm

Graphite proportion , %

Carbon percentage dissolved imposed , %

Soaking time at sintering , min

1050 60 1 0,4 35,8

1050 60 1 0,6 58,4

1050 60 1 0,8 87,6

4. The experimental check of the mathematical model

For the experimental check of the proposed algorithm , has been realised the powder mixture sintering formed 99 % Fe and 1% C(gr) , at the temperature of 1050 oC , with different soaking time: 30, 60, 90 minutes. After the sintering, it has been realised the microscopic analysis. The micrografics are presented in figures 4 – 5 and show a good conformity between the calculated values and the experimental ones of the C percentage dissolved in iron.

Fig. 4 Material Fe + 1 % C. The soaking sintering time: 30 minutes

Atac Nital 4% ( x 200)

Fig. 5 Material Fe + 1 % C. The soaking sintering

time: 60 minutes Atac Nital 4% ( x 200)

Fig. 6 Material Fe + 1 % C. The soaking sintering time: 90 minutes

Atac Nital 4% ( x 200)



66

5. Conclusions

The presented thermodynamic aspects allows the estimation of the mechanisms through which is realised the carbon dissolvation into iron during the sintering : - at temperatures up about 800 °C through the CO dissociation into C in “birth” state and CO2. - at high temperatures through the C volume diffusion into iron particles .

The simulation of the sintering iron-graphite powders mixture allows the soaking time determination for obtaining sintered pieces, with maximum resistance, with a C content of 0,7-0,8% correspondent to the eutectoide composition. .

Is approached a computational integration method of the diffusion equations, using the finite differences method. The written programme using the presented algorithm calculates the soaking sintering time depending on the temperature, on the powder granulation, on the graphite proportion and on the C percentage value dissolved into iron imposed by the user.

The metallographic analysis realised using sintered tests with different soaking time shows a good conformity between the calculated values and the experimental ones of the C percentage dissolved into iron.

When also the copper is present in the antifriction material, for the use of the programme, which moulds the sintering, are needed corrections to the C adsorption speed constant at the grains surface and to the C diffusion activation energy into austenite.

References [1]. Oprea, F., 1988, “Teoria proceselor metalurgice”, Editura Didactica si Pedagogica, Bucuresti [2]. Cojocaru, M., 1998, “Modelarea interactiunilor fizico-chimice ale produselor metalice cu mediile”, Editura Matrixrom, Bucuresti. [3]. Petrea, I., 2003, “Cercetari privind obtinerea si proprietatile trobologice ale materialelor compozite poroase in matrice de fier”, Ph.D. Thesis, University of Galaţi. [4]. Williams, J. A., 1994, “Engineering Tribology”, Oxford University Press, Oxford. [5]. Surdeanu, T., Pernes, M., 1984, “Piese sinterizate din pulberi metalice”, Editura Tehnica, Bucuresti



67

CROSS – SECTION MODIFICATION, PENDING

OPEN DIE FORGING PROCESS

Doru C. HANGANU

“Dunărea de Jos” University of GALATI e-mail:[email protected]

ABSTRACT

This paper presents same experimental tests, to determine cross – section

modifications of cylindrical work piece, pending forge drawing process. To work done this paper is used theoretical knowledge and experimental testing, performing on lead cylindrical samples, about cinematic flow of metal.

KEYWORDS: algorithms, simulation, forge drawing process

1. Introduction

Free Forge drawing is an important operation in free forge case, special for get forging kind arbors. The flow chart of forge drawing is shown in figure no.1 Pending forge drawing the material is flowing in two directions: longitudinal and lateral, according to less deformation resistance low: kinematics flow material is proportional with polygon’s area who is obtained on the frontal surface, like figure 2.

This is conditional from width of deformation tool.

Friction forces occurrence involve maker the cross-section to have barreling, like figure 3.

2. Work mode presentation

For simulation the material flow pending forge drawing process for cylindrical work piece, between plane tools, was used plumbs cylindrical samples, with Ø50 x 60 mm, who was plastic deformation on hydraulic press.

Fig.1 Flow chart of forge drawing process

Upper smith anvil

Semi product billet

Lower smith anvil

B

Deformation was making with following length clamp up: 40 mm, 50 mm and 60 mm. Then, after plastic deformation with different deformations degree, was measured the following data:

• H1 is Height after deformation; • Dmin – Minimum diameter; • Dmax – Maximum diameter.

Dmin

Dmax

H1

Fig.3 Measured dimension data after deformation

Fig.2 Flow material according to less deformation resistance low




68

In this cross-section there are a part without barreling A1 and a part with barreling, A2 like figure 4.

The barreling in this case is a circle barreling.

A1

A2

H1

Dmin

Fig.4 Cross-section after deformation

First calculate the volume compass

between the deformation tools, with relation: BRV ⋅⋅= 2

0 π After deformation we can observe that two parts of cross section is barreling A2 (hatching parts) and a part without barreling A1. Cross surface area may be calculate like sum: 21 2 AAS ⋅+= The calculus mode of cross section, but special for calculus circular segments area, may be observed in figure 5 shown below: Because A1 surface is a rectangle, its area on calculate with relation: 1min1 HDA ⋅= Geometrical form of A2 surfaces is two circular segments. In this case we known the span AB = H1, without known radius OA. trcs AAA −=2

Where: Acs is circular segment area; Atr – AOB triangle area.

ϕπ

⋅⋅

=360

21RAcs

Where: ϕ is the AOB central angle

2

)( 11 fRHAtr−⋅

= ;

Where: f is highness of circular segment (hatching). R1 – barreling radius The highness of circular segment is calculate with relation:

22minmax DD

f −= ;

( )4

212

12

1HfRR +−= ;

f

HfR⋅

+=82

21

1

( )fRHtg−

=1

1

22ϕ

;

( )

236011

21

2fRHRA −

−⋅⋅

=ϕπ

;

Knowing cross section area after deformation and volume before deformation, may be calculate length after deformation B

φ O

A

B

Fig.5 Calculus of A2 area

f R1

H1

B1, with relation:

21

01 2 AA

VB

⋅+= ;

The deformation mode is shown below on the cylindrical work pieces, figure 6 - 12.

Fig.6 Initial cylindrical sample



69

Fig 7 Cylindrical sample after deformation on 50 mm capture length Fig.10 Work piece pressed on integer

initial length, with high deformation degree

Fig.8 Cross section of deformation zone

Fig.11 High view of cylindrical sample after deformation with 25% deformation

degree

Fig 9 Work piece pressed on integer initial length, with low deformation

degree

Fig 12 High view of cylindrical sample after deformation with 70% deformation degree and

40 mm capture length.

4. References [1] V. Chiriţă, I. Drăgan, Al. Maniu, A. Vasiliu – Matriţarea la cald a metalelor şi aliajelor, Editura tehnică, Bucureşti, 1979; [2] I. Dore Landau – Identificarea şi comanda sistemelor, Editura Tehnică, Bucureşti, 1997; [3] M. Ghinea, V. Fireţeanu – MATLAB – calcul numeric- grafică – aplicaţii, Editura Teora, Bucureşti, 1998

]4] V. Popescu – Forjarea şi extruziunea metalelor şi aliajelor, Editura didactică şi pedagogică, Bucureşti, 1976. [5]. V. Moldovan, A. Maniu – Utilaje pentru deformări plastice, Editura didactică şi pedagogică, Bucureşti, 1982 [6] E. Cazimirovici – Bazele teoretice ale deformării plastice, Editura BREN, Bucureşti, 1999 [7]. M. Stan – Mecanica mediilor continue – Editura Matrix Rom, Bucureşti, 2001



70

NONFERROUS ALLOYS WITH SPECIAL PROPERTIES

IN HIGH TEMPERATURE

Maria VLAD, Olga MITOŞERIU, Emil STRĂTULAT

“Dunarea de Jos” University of Galaţi e-mail: [email protected]

ABSTRACT

Casting forms materials used in both ferrous and non-ferrous alloys continuous casting must be hard, wear resistant and very good thermal conductivity.

Our researches had in view to replace Ag with other metals (Cr, Zr) with a view to reducing costs of continuous casting and obtaining cheaper materials of high durability and thermal properties close to those of the Cu - Ag alloy. In order to improve mould durability as well as to get high-quality machine ports, the metallic moulds must be covered with nonmetallic, metallic and metaloceramic coatings. Experimental researches have tried to obtain thicker (about 300 μm) and hard layers.

KEYWORDS: continuous casting, mould, nonmetallic, metallic and

metaloceramic coatings, durability and thermal properties

1. Influence of small alloying elements on the physico-mechanical characteristics of

Copper alloys used in metallic moulds Nowadays, both in our country and abroad, the crystallizers used in steel continuous casting [2] are made up of Cu or Cu-Ag. Pure copper is used only for small crystallizers, because within the 140 - 180°C range when it reduces its hardness to half . Anyway it is not used in big crystallizers as it undergoes structure and volume changes, namely it contracts because of unfavorable yield properties [3]. But copper properties can be improved by alloying and cold deformation methods . Analyzing the influence of small Cr or Cr-Zr additions at room temperature as compared with the Cu-Ag alloy, we found that, according to Table 1, both the refractoriness and the mechanical properties are improved which influence, to a certain extend, the

thermal conductivity or the thermally untreated materials.

Chromium and zirconium, heavily fusible materials, increase both the melting and re-crystallizing temperature of copper, making up new phases with complex structures. The structure of the 0.80 %Cr alloy is given in Figure 1 where the solid solution α of chrome in copper and un-uniformly distributed chrome crystals are shown. The zirconium additions in Cu-Cr lead to the improvement of the mechanical properties at temperatures of about 400°C, due to CrZr2 and Cr2 Zr compounds which are formed, Figure 2. These alloys are structurally hardened by annealing thermal treatments at 400-500°C, tempering in solution at 1,000-1,500°C and quenching within 450-500°C. The physical-mechanical properties at 400°C after thermal treatments are shown in Table 2.

Table 1. Physico-mechanical properties of studied copper alloys

Chemical composition (%) Physical-mechanical properties Alloy type Cu Ag Cr Zr λ (W/m⋅K) HB Rm (daN/mm2)

Cu-Cr rest - 0,8 - 103 66 24 Cu-Cr-Zr rest - 0,7 0,2 104 69 28

Cu-Ag rest 0,08 - - 496 40 19




71

Fig. 1 Cu-Cr alloy (0,80% Cr) microstructure thermale untreated.

Fig. 2 Cu-Cr-Zr alloy (with 0,7% Cr and 0,2% Zr) thermale untreated

Table 2. Physico-mechanical properties after thermal treatments

Chemical composition

(%) Physico-mechanical properties Alloy type

Cu Cr Zr Ke (Ω-1⋅m-1) λ (W/m⋅K) Tsof (K) HB R (daN/mm2)

Cu-Cr rest 0,8 - 27,5⋅106 414 748 122 32 Cu-Cr-Zr rest 0,7 0,2 26,9⋅106 405 773 134 34

Thermal conductivity may be calculated by means of Wiedemman-Franz-Lorenz relation:

TkL

e

λ=

where: ke-thermal conductivity , L-Lorenz constant (L=2,22⋅10-8) The improvement of all properties is shown to be the result of an unstable homogeneous and supersaturated state after tempering that makes the material to be relatively soft quite like normal copper, the tempering treatment took place at 450°C when the chrome excess precipitated from the solid solution, while the compounds have uniformly been dispersed in the base mass (solid solution α).

2. Researches concerning the increase of crystallizers durability in the steel

continuous casting by electrolytic metallic deposition

Crystallizers used in steel continuous casting are made up of pure copper or chrome alloys and therefore have a small wear resistance with entails the appearance of flaws in the crystallizer wall. In order to increase durability in all these shapes metallic

deposition with chrome, nickel, nickel alloys or multilayer protection can be used, Figure 3.

Fig. 3. Multilayer deposition In other works we are show that for thick nickel layers (1-3 mm), the copper plate durability can improve up to 4,000 chargers versus 250. To obtain thicker layers several variants of Watts electrolyte solutions have been analyzed to whom a series of organic and un-organic chemical substances have been added with a view to obtaining the required quality. For all variants optimum parameters which determine them have hydrogen ion discharge on the cathode height of deposited layer. The use of electrolyte solutions without additions [4] induced pores, pittings Figure 4 due to inclusion of hydrogen emitted as well as to the colloidal hydroxide absorbed on the cathode.



72

Fig. 4 Electrolyte deposition with nickel on copper with pittings, pores and inclusions Fig. 5 Thick and hard nickel coating on copper

Electrolyte composition sensibly affects the

hardness of cathodic depositions. Additions of ammonia salts, fluoride, chloride, etc in solutions induce a rise in hardness. For example, during nickel deposition from an ammonia sulfate solution, microhardness has risen from 120 HV to 250 HV, while the presence of a small quantity of chlorides increased microhardness to 320 HV. The excess

ions must be avoided which at a higher level, leads to a nickel and ammonia double sulfate formation which can be separated and make up macroscopic impurities. The influence of current density on cathodic efficiency has been analyzed for a Watts electrolytic bath (matte nickel) and of a hard-deposition electrolyte Figure 6.

+4NH

Fig. 6. Current densyty influnce upon cathodic efficiency.

It is very important to notice the influence of

varying pH between 2 and 5.5 upon the current catholic rate and hardness in nickel coating. The hardness of galvanic depositions depends on their structure, the finer of the structure. For getting thick layers, special measures must be taken in as far as surface preparation is concerned: lack of asperities, detensioning to 200-220°C and slow cooling, pore, nclusion and flaw elimination. i

3. Conclusions

By microalloying copper or copper alloys with heavily-fusible metals, rare metals (Cr, Zr) their mechanical properties and refractoriness can be improved without influencing thermal and electrical

conductivity too much. Big durability of copper or copper alloys metallic forms at high temperature (500-800°C) can be achieved by using thick and hard metallic layers with chrome, nickel or their alloys. Better results are obtained during nickel coating, because nickel thermal dilatation coefficient is almost qual to that of copper. e

References

[1]. Hans, Schrewe, : Stranggiessen Von Stahl, 1987, RFG [2]. Laurentiu, Sofronie, : Turnarea prin cadere libera in forme metalice, Editura stintifica si tehnica, 50-85, 1983 [3]. Frederick, Lowenheim, : Modern Electroplating, 10-200, 1974 - U.S.A. [4]. Oniciu, L., Grunwald, F. : Galvanotehnics, Ed stintifica si

tehnica,17-23, 1980



73

MATHEMATICAL MODELING OF ROTATION ÎN CASE OF VERTICAL CENTRIFUGAL CASTING

Aurel CIUREA; Marian BORDEI, Dorin ETIMIE


ABSTRACT

Albeit the principle of centrifugal casting is known since 1816 (Germany), this

method found its practical application a century later. One of the issues was represented by correct determination of facility rotation for centrifugal casting. The present work refers to a mathematical pattern for rotation determination in case of vertical centrifugal casting.

KEYWORDS: pearlitic structure, free cementite, soaking time at sintering,

antifriction material

1. General considerations

By seeing a liquid rotation in a vessel, we find that, under the action of centrifugal force, it trends to go to those places of the vessel located at the biggest distance from its rotation axis. The distance between liquid particles and rotation axis depends on the rotation of the vessel. The more the rotation grows, the more the gravity force is going down and the liquid is lifting on the inner walls of the vessel. At a certain rotation, big enough, it goes out through the upper outlet of vessel. A constant rotation leads to a liquid layer having the form of a paraboloid. That is used in centrifugal casting, where the liquid is limited towards the enter wall of the

vessel under rotation, while inner limit is imposed by the centrifugal force under action.

In its turn, the centrifugal force depends on rotation and must act upon the metal till the later is completely solid.

In considerations below, the liquid metal will be considered as an ideal liquid.

2. Mathematic deduction of rotation equation

We will have in view the representation of a centrifugal rotation formula necessary for a cylindrical object, empty inside, depending on: radius, wall thickness and its height. For the beginning we take into consideration the area of a

· ·

Fig. 1 The diagram of vertical centrifugal casting factor




74

liquid under rotation, fig. 1. A radial force is acting upon friction of

liquid particles (centrifugal force), Fc and a vertical force (gravity force), G, directed downside.

(1) 2c rmF ω⋅⋅=

G =m g (2) where:

m - the liquid mass; r - radius; g - gravity acceleration; The force R, as a result of Fc and G, acts in a

certain point and is perpendicular to a tangential plan of this point.

The following relation is to be found in the forces parallelogram:

ϕ=ω⋅ tg

mgmx 2

(3)

The differential quotient representing a measure of tangent inclination is as follows:

ϕ= tgdxdy (4)

From equation (3) and (4) we get the following:

(5) dygdxx 2 ⋅=⋅ω⋅

By integration we get the equation of a parabola:

∫ ∫=⋅⋅ω dygdxx2

or:

yg2

x 22⋅=

⋅ω (6)

By changing, we get:

g2

xy22 ω⋅

= (7)

where:

T2π

=ω the circular frequency with n60T =

where: T - the duration of one rotation; n - number of r.p.m. By introducing the above mentioned values

we get in (7):

g1800xn

3600g2nx4x

Tg24y

222

2222

2

2

⋅⋅⋅π

=

=⋅

⋅⋅π=⋅

⋅

π=

g1800xny

222

⋅⋅⋅π

= (8)

The form of the liquid metal gets in centrifugal casting inside a cylindrical object is represented in fig. 2. We can see that the angle α and the inner form of the liquid depend on rotation for the given radius and length of the object.

At bigger rotations (>100) the parabola will have a slight curve which can be approximately represented by a spring going through points P1 and P2 located in the parabola. were:

a - the wall thickness; x0 - the radius of centrifuged object;

Fig. 2 The parabola formed during the process of centrifugal casting inside an empty cylindrical object (vessel)


FASCICLE IX METALLURGY AND MATERIAL SCIENCE, ISSN 1453-083X NR.1 2004

yo - the height;

10x o - effective thickness of object ‘s wall at

the lowest point. If we note in equation (8) the following:

g1800

nC22

⋅π

=

it results that:

Cyx

xCy 2

=

⋅= (9)

By the help of above equation we can establish the coordinates of P1 and P2.

( ) ( ) ( )[ ]200111 axC;axy;xp −−= (10)

( )

( ) ( )[ ]⎪⎭

⎪⎬⎫

⎪⎩

⎪⎨⎧

−−−−

=

=

02

00

20

222

yaxC;C

yaxC

y;xp

(11)

The inclination of parabola spring corresponds to the tangent we are looking for according to fig.2:

α=−

tga

10x

y

0

0 (12)

In this equation, the distance between the enter wall of the cast object and imaginary line of parabola, was considered for sockets with an average

value 10x0 . It can be It can be replaced by any other

value: From fig.2 it results:

( )a10xy10

a10x

ytg

xxyy

0

0

0

0

21

21−⋅

=−

=α=−−

(13)

By replacing the values for x1, x2, y1, y2 (13)

it results:

( ) ( )[ ]( ) ( )a10x

y10

CyaxC

ax

yaxCaxC

0

0

02

00

002

0−⋅

=−−

−−

−−−−

( ) ( )a10xy10

CyaxC

ax

y

0

0

02

00

0−⋅

=−−

−−

(14)

By solving the equation (14) it results:

( )

CyaxC

910x 0

20

0−−

=

or:

( )

CyaxC

81100x 0

202

0−−

⋅=

where we get:

( ) ρ⋅π

=+−

⋅=

1800n

a100ax200x19

y100C

22

20

20

0 (15)

From this equation it can be calculated the theoretic rotation for centrifugal casting:

( )

20

20

0

220

20

0

a100ax200x19

y4230n

a100ax200x19

g1800y100n

+−⋅=

π+−

⋅=

[ ] (16) 12

02

0

0 minaax2x19,0

y423n −

+−⋅=

3. Establishment of useful rotation

The above equation is valid only for an ideal liquid. The materials used in casting are different so that determination of rotation must be connected with the liquid condition under casting. Also, it must be taken into consideration the characteristics of casting shape, the way the liquid metal is introduced into the shape and the cooling conditions.

The material liquid conditions which is under casting must depend on many factors, namely:

- material compactness; - viscosity and friction in casing shape;

Such factors, depending on respective specific conditions, are very difficult to be theoretically determined, that is why they can be included under the form of a constant K0. The latter is experimentally determined by centrifuging a part at various rotations.

As a result the required rotation is determined by the following relation:

[ 12

00

00 min

aax2x19,0

y423Kn −

+−⋅= ] (17)

75



76

By introducing these values into the equation (16) we get the theoretic rotation: When determining the constant K0 it must be

chosen the rotation to which a correct configuration of cast piece is obtained on one hand and a material with proper characteristic on the other hand. For serial production, casting condition be observed when determining the constant K0 and be observed first of all, maintaining the casting temperature.

[ ]1min125416,0222519,0

8423n −=+⋅−⋅

=

Experimental casting was made on a profiled

bonze allay, the tests being made for the rotation as fallows:

When establishing the welting temperature, it must be taken into account the heat lass by cooling during liquid metal transfer from the melting pot to the centrifugal facility. The way the shape is supplied with metal chosen when casting, must be maintained in order to avoid unpleasant surprises during production.

a) 1254 rot/min => K0 = 1 b) 1300 rot/min => K0 = 1,04 c) 1350 rot/min => K0 = 1,076 d) 1400 rot/min => K0 = 1,116 The centrifugal rotation is not therefore the

only factor important to the quality of cast piece. All above mentioned factors must be taken into account, for example aluminium has a strong tendency to absorb gases.

The best results were got for n = 1350 r. p.

m, K0 = 1,076.

5. Conclusions

The equation of rotation theoretically calculated is valid for an ideal liquid, for practical cases being used as starting point for preliminary tests of technology for centrifugal casting.

Centrifugal casting represents a technical method by witch there are obtained pieces with most compact structure compared to shell casting and casting under pressure. If we take the conductivity as compactness measure unit for a cast piece, we get the following different characteristics:

4. Examples of practical use rotation equation

Fig. 3 represents an example to show the way this equation is used

Fig.3. Centrifugal cast bresh.

a) Cast pieces under pressure, conductivity of maximum 31 Ω mm2/m;

b) Shell cast pieces, conductivity of maximum 34 Ω mm2/m;

c) Centrifugal cast pieces, maximum 35 Ω mm2/m.

d) Centrifugal cast pieces are characterized by very big compactness, a fact which can be considered of a great importance for many special pieces.

References [1]. HORST – GŰNTER HALT – Turnarea centrifugală a

metalelor uşoare – Bucureşti – 1960; The bearing socket in fig. 3must have the length y0 = 8 cm, radius x0 = 5 m and wall thickness a = 0,4 cm.

[2]. D. TALOI Ş.A. – Optimizarea proceselor metalurgice – E.D.P., Bucureşti – 1983.



77

THE HARDENING OF THIN BANDS BY PLAIN CARBON STEEL

WITH LOW CONTENT OF CARBON BY QUENCHING

Petrică ALEXANDRU

“Dunărea de Jos” University of Galaţi [email protected]

ABSTRACT

For the plain carbon steels with very low carbon content the usual method for

improve the strength is the cold work hardening and the quenching is considerate in industrial practice an idle technology for this category of alloys. The researches was orientated to the obtained by quenching for steel strapping for packaging. The quenching was applied with three hardening baths: water, salt-water (5% NaCl) and salt-water (10% NaCl). The test data of quenching steel strapping approved the value of this technology by compare with cold work hardening.

KEYWORDS: plain carbon steels, very low carbon content, steel strapping,

quenching.

1. Introduction

The high values for the mechanical characteristics (the breaking strength, yield resistance) is done, at large, in the material metallic case, through two category of conventional proceedings: cold-working and the thermal treatment.

For material metallic that don't suffers the capable solid his transformations or if these transformations don't influence significantly the marked features, the cold hardening as the effect of plastic deformation to cold is the unique option for increased the mechanical characteristics. For the plain steels with low content of carbon, in the industrial practice is the tendency to consider as if are not significantly hardening by quenching.

Many specialists ignore the fact that for thin semiproducts can achieve by cooling, from the austenitic domain, a sufficient cooling rate for the transformation austenite-martensite let us can be in progress.

In this article are presented the experiments of hardening by thermal treatment from the plain steel carbon bands of packaging with the section 0,5x10-mm mark A3k. It was selected a especial charge with low content of carbon, 0,04%, just tested to limit, the possibility to obtain through thermal treatment the mechanical characteristics that is enforced for this bands. Be due to assure a very high tensile-strength, but and a breaking elongation of precinct 4-5%, as good as a feast assured the resistance to the shocks which appear to manipulation or transport of heavy packs.

0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,8 0,90Carbon [%]

Tem

pera

ture

[0 C]

700

600

500

400

200

300

800

Fig.1. The start temperature of transformationaustenite-martensite variation, Ms, with carbon

content.

2. The technological characteristics of thermal treatment for used steel

With the growth contents of carbon,

temperature whereat it begins the martensitic transformation is displaced to low values, just as accrue from the figure 1. For contents extremely low in carbon (thousandth by percent) temperature Ms= 7500C, and for the usual area contents, in the steel bands used for deep-drawing, Ms= 530-5500C.

Very high temperature whereat begins the martensitic transformation permits the carbon separation from martensitic matrix as the effect of self-tempering. The amplitude of self-tempering




78

depends great deal of cooling speed. Thereby, result take over from literature, as regards the mechanical characteristics dependency for the soft steels of the content of carbon, must be correlate with the thickness products and with the properties of quenching medium used for hardening. The difference by chemical composition, conditions of cooling (mode of agitation, for same medium of quenching), etc., can be responsible for the waged concordance of the literature dates. The literature offers a relatively reduced number of diagrams CCT, for the steels with reduced content of carbon. The

concordance of the date contained in these diagrams given from different sources is insufficiently. Thus in diagram CCT from figure 2, result that Ms= 4800C for the steel with 0,06% C, as what from figure 1 result for some contained of carbon Ms= 5400C. It is enforced, failing of a precise characterization, from angle hardenability steels with very low content of carbon, as through experiment let us is caused the conditions of thermal treatment, which permit the procurement desirable properties. .

0,1 10 102 103 1041

300

500

700

Ac3=9000C

Ac1=7200C

Ms=4800C

%C=0,06

Time [s]

Tem

pera

ture

[0 C]

V1V2

F

P

A

Ms

M

V1=104 0C/sV2=300 0C/s

Fig. 2. CCT diagram for plain carbon steel with 0,06 %C.

2. Experimental conditions 2.1. Objectives

The purpose it was determination of impact by the cooling capacity quenching medium of the mechanical characteristics for the steel strapping with low content of carbon by thermal treatments. It was select a steel with reduced content of carbon, starting from premise that for a such hardened steel is obtained a sufficient plasticity, just without the adhibition treatments of tempering.

It tried the show-down of possibility to obtain the bands of packing with superior performance, through thermal treatment for steels with reduced content of carbon, who are considerate as a rule, as be practical non-hardening alloys by quenching.

2.2. Samples The experimental-use samples were bands with

the transversal section 0,5x10mm and length of

650mm. Material was the steel for deep drawing, obtained by cold rolling, A3k. The chemical compliant composition (STAS 9485-80) and the proprieties able to mechanical delivery (reannealed) are presented in table 1 and 2.

Table 1

Mechanical characteristics of A3k steel

Rp0,2 [N/mm²] max. 240 Rm [N/mm²] 270÷370

A5 [%] min. 31

Must underline, online the chemical composition, as it followed the of a use steel with a very low content of carbon. Thus, from the reasons presented was selected, on the strength of analysis notes from manufacturer (ISPAT-SIDEX S.A., Galatzi) a steel with 0,04%C.

Table 2

Chemical composition [%] Mark

C Mn Si P S Cr Ni Cu

A3k max. 0,08

0,20÷0,4

max. 0,03

max. 0,03

max. 0,035

max. 0,06

max. 0,08

max. 0,08



79

2.3. The experimental installation of quenching

In order to succeed the heating and cooling

with high rate of thin bands it conceived a special installation, presented in figure 3, for achieved these objectives.

The heating is done in a electric furnace (1), the band (8) it is in position to a muffle from refractory steel (9) that assure the protection to oxidation, maintain of a rectilinear form of band and guide the band in period of transfer to quenching medium (6) find out in the tube (7).

1 2 3 4

5678910

Fig. 3. The experimental laboratory stand used for the hardening of strapping steels,0,5x10x650 mm by plain carbon steel.

1-electric furnace; 2-electric resistor; 3-steel wire; 4-traction spring; 5-mounting support; 6-hardeningbath; 7-steel pipe; 8-steel band; 9-steel refractory muffle; 10-thermo-element.

The hardening is in progress through the

liberation spring (4), what run-down and by wire of steel (3) draw the band from furnace to medium of quenching, with very high speed.

2.4. The quenching mediums

Three types of quenching mediums, with temperature 100C, were used:

1. water; 2. salted water with 5% NaCl; 3. salted water with 10% NaCl.

3. Results

Micrographies of hardened steels presented in figures 4-7 relived the acicular relative fine appearance of structure, composed by supersaturated ferrite in carbon and martensite. The austenite

transformation, with very high speed, leads to a big number of structure defects, type dislocations in main, and what to their row causes hardening of steels.

By analyze of the mechanical characteristics of hardened band result:

1. The tensile strength with the eldest value (1030 N/mm2) is obtained in the case of quenching medium (3) used for cooling (watery solution with 10% NaCl).

2. The elongation with the maxim value is obtained in the case of quenching medium (2) used for cooling, (watery solution with 5% NaCl).

3. The deformation energy by traction to deformation rupture, is maximum in the case of quenching medium (2) used for cooling, (watery solution with 5% NaCl) and minimum in case of quenching medium (3) (watery solution with 10% NaCl).



80

Fig. 6. Microstructure of hardened band,cooling medium (2), (x300)



Fig. 4. Microstructure of annealed band,(x300).

0

200

400

600

800

1000

1200

1 2 3

Hardening bath

Rm

[N/m

m2 ]

Fig. 8. The tensile strength variation with type of quenching medium.



81

0

1

2

3

4

5

1 2 3

Hardening bath

A10

[%]

Fig. 10. The deformation energy by traction to deformation rupture variation with type of hardening bath

4. Conclusions

The steel strapping with low content of carbon (approximately 0,04%) and thickness of approximate 0,5 mm can touch after hardening by thermal treatment a value of breaking strength by 1000 N/mm2 and a breaking elongation by 4%. These represents the very good performances in the domain of bands by packing, because currently values obtained by cold-working are maximum 750-800

N/mm2, for breaking strength and 1-2% for breaking elongation

References [1]. K. Günther, Zur Festigkeitssteigerung Kohlenstoffarmer unleigerter Bänder durch Wärmbehandlung, Neue Hütte, nr. 12, 1969, pag. 734-739. [2]. Schrader, A., Rose, A., De Ferri Metalographia., Verlag Stahleisen m.b.H., Dűsseldorf, 1966.



82

STUDIES AND RESEARCHES TO IMPROVE

THE STRUCTURE AND HARDNESS OF STEEL C420

Stela CONSTANTINESCU, Olga MITOŞERIU

“Dunărea de Jos” University of Galati Stela.Constantinescu@ email.ro

ABSTRACT

The present paper attempts to analyze the results obtained after applying a

number of 4 heattreatments in order to fiind those results which best fit the plate qualityconditions . The results of the experimental heat treatments met the requirements for microstructures. The microstructures show , apart from the chemical non- homogeneity and non-homogeneity as regards the real grain size , which makes the final structure be quite different even with the same sheet at the head and end of the plate and between the surface and core of the plate . Measuring the hardness of across the samples was made by Brinell method . The influence of annealing temperature is normal both in terms of surface hardness and core , and its value decreases while the tempering temperature increases .

KEYWORDS: tempering temperature, value of hardness, heat treatment, microstructures

1. Introduction

Since the mechanical properties of any plate are

directly dependent on the final structurereached after applying quenching and annealing treatments , the present paper attempts to analyze the results obtained after applying a number of 4 heattreatments in order to fiind those results which best fit the plate qualityconditions .

For this purpose , during the whole manufacturing flow (fig.1), the authors watched a

charge along its entire technological , for thick plates made from alloyed steel of high yielding point usually employed for robust constructions of high breaking strength such as : fixed or moving store tanks , rolling paths , electrical line pillars .

The steel is alloyed with Mn and belongs to the category of fine granulation highyielding point steels . The chemical composition as determined on charge is at table no. 1

Prod

uctio

n, %

Years

Fig.1.The product of steel C420



83

Table 1. Chemical composition of steel C420 Steel

Chemical composition

grade C Mn Si S P Al As Cr Ni Cu V Mo C 420 liquid

0,15 1,24 0,17 0,01 0,019 0,03 0,007 0,09 0,22 0,013 0,02 0,02

C 420 product

0,17 1,27 0,15 0,01 0,019 0,03 0,007 0;09 0,22 0,013 0.02 0,02

For a deeper insight the steel concerned , the fabrication and rolling conditionswere first studied . To avoid the noxious effect of the residues : P , Al , S , C , the steel was made in an electrical furnace of 60 t in SIDEX S.A. with or without degassing inside thecorresponding vacuum devices . A number of 4 ingots of 15 t were obtained

Cechiv. = C +

6Mn +

5VMoCr ++ +

15CuNi +

Cechiv. = 0,15 + 1 24

6,

+

0 09 0 02 0 025

, , ,+ + +

0 22 0 1315

, ,+ (1)

Cechiv. = 0,3981 % It can be noticed that the experimental steel meets the required technical conditions (fig.2.) The brams were heated in propulsion furnaces according to the scheme below in table no.2

.

Thermal effect Carbones echivalente [KJ/cm] [%]

15

CuNi5

VMoCr6

MoC ++

++++

Fig.2.Thermal effect and carbones echivalente

Table 2 . Propulsion furnace temperature at o C Preheating area Heating area Equalizing area

1030o C - 1060o C 1240o C - 1270o C 1210o C - 1230o C

2. Experimental researches For the purpose of lab rearches samples were cut from the plates.

The critical points were experimentally determined by a derivatograph and checking was performed with the R.A.Grange relations . AC1 = 723 -14(Mn + Ni) + 22(Si + Cr)=722o C (2)



84

AC3 = 854 - 180 C - 14 Mn - 18 Ni + 45 Si + 1,7 Cr = 814o C The accurate annealing temperature is egual to : TC = AC3 + (30 - 50)o C = 830o C + (30 - 50)o C (3) Heating for superheating austenitic phase was carried aut , this is possible because the steel has a fine granulation and therefore there is no danger for the austenitegrains to grow exaggeratedly . The following 4 annealing regimes were established , each of them followed by tempering to 3 different temperatures : A. Annealing temperature 880o C, water cooling, tempering temperature: 620oC; 650oC; 680oC; Water cooling. B. Annealing temperature 925o C, water cooling, tempering temperature : 620oC; 650oC; 680oC; Watter cooling . C. Annealing temperature 950o C, water cooling, tempering temperature: 620oC; 650oC; 680oC; Water cooling .. D. Annealing temperature 980o C, water cooling, tempering temperature,: 620oC; 650oC; 680oC; Water cooling . Experiments were carried aut on samples of : 250 x 120 x 40 heated in furnace with bars of type AUTOMATICA. For austenitization purpose :

τinc = K W = K )hlblbL(2

bhl++

τinc. = 40 )4x25()4x12()25x21(2

4x25x12++

= 536 min. ( 4 )

τtot. = K1 x τinc. = 1,4 x 53,6 = 75 min. (5) The results of treatments applied have been assessed according to the metallographic structure into the plate thickness on a cross section perpendicular to the rolling direction and by measuring the hardness on the same cross section .

3. Results and conclusios

The results of the experimental heat treatments met the requirements for microstructures . The study of these microstructures show that the steel undergoes a strong segregation phenomenon both along the plate and across which further leads , after applying the same treatment , to different structures at the end , head and section of the plate . A sort of chemical non homogeneity is found with respect to plate end head and also between the surface and core , the different contents of C in the plate thickness makes the areas of low C content unable to reach the annealing structure or makes it impossible the formation of relatively rough carbons . The microstructures show , apart from the chemical non- homogeneity and non-homogeneity as regards the real grain size , which makes the final structure be quite different even with the same sheet at the head and end of the plate and between the surface and core of the plate . It can also be seen a significant amount of inclusions of different types of non uniform distribution and size which negatively affects the values of the mechanical properties of steels . Measuring the hardness of across the samples was made by Brinell method , and a number of min. 3 determinations were obtained for each sample , the average value of these measurements is given in table no.3 .

.Table 3 . Measuring the hardness

Mean value of hardness, HB Regimes of Tempering temperature o C treatment 620 650 680 Remarks

SI M SII SI M SII SI M SII A 211 184 219 198 181 198 191 180 187 Top

T = 880o C 211 215 226 215 198 198 191 180 101 bottom B 191 184 187 191 177 187 180 167 180 Top

T = 925o C 198 195 187 198 187 191 188 180 191 bottom C 229 221 229 229 209 220 215 191 202 Top

T = 950o C 234 224 234 189 184 198 190 187 174 bottom D 211 191 211 202 187 190 174 158 179 Top

T = 980o C 234 239 239 219 219 234 187 180 180 bottom



85

It can be seen that the hardness values increase with higher austenite temperatures . The influence of annealing temperature is normal both in terms of surface hardness and core , and its value decreases while the tempering temperature increases . From the foregoing it follows that the C and D regimes provide more uniform results . It is worth mentioniing that the presence of a large number of inclusions and chemical non-homogeneity affect the mechanical properties of plate irrespective of the steel structural condition .

References [1]. R. Bruscate: Temper embrittement and acep embrittement of 21/4 Cr-Mo schielded metal Are Weld Deposits Welding Journal, aprilie 1970 . [2]. Walanbe T.: Mechanical proprieties of Cr-Mo steels after eleveted temperature service, Partea I: Document II S-IX-116-79; Partea II: Document II S-IX-1167- 1990 [3]. N. Steven: Jise, vol.193, pag.141, 1989 . [4]. Sugiyama T. : Kobe Seel Engineering Reports, vol.25, nr.4, 1995. [5]. Delman B. - Introduction a la cinetique heterogene, cap.V, VI, VII, VIII, IX, Paris, 1989.

[6]. Kelly R.G., Newman R.C. - Electrochemical and Optical Techniques for the Study and Monitoring of Metallic Corrosion”, Klwer Academic Publishers, Series Applied Sciences, Vol.203/ 1991, pag.665-695. [7]. Knauschner A. - Elektrochemische Metallabscheidung, Oberflachenveredeln und Planttierent von Mettalen, Leipzig, 1988, pag.61-119. [8]. Ralph E., Bockins J., Conway B.E. - Modern Aspects of Electrochemistry, New York and London, nr.21/1994 , p.29-58. [9]. Constantinescu S. – Studies on Thin Carbide and Nitride Layers Deposition on Metal Basis, Based on Chemical Reaction at High Temperatures: Academic Thesis, Galati , 1998, p. 84 . [10]. Constantinescu S., Drugescu E.- Studies on Mechanism and Kinetics of Phase Transformation in Superficial Layers Using Unconventional Procedures. Research Contract no. 5005, Galati ,1995, p. 53. [11]. Constantinescu S., Drugescu E.,Radu T. - Practical application of AE of the different grades of steel. “ Proceedings of the 25thEuropean Conference on Acoustion Emission Testing “ EWGAE 2002 “ Prague Czech Republic , 11 –13 september, 2002 , p. 135. [12]. Constantinescu S. - Influence of manufacturing process on chemical and structural homogeneity of welded pipe sheets for tanks and vessels working under pressure. “ Proceeding of the International Conference on Advances in Materials and Processing Technologies, september 18 – 21 , 2001, Leganes, Madrid, Spain, p.57 .

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