CONTENTS

SCIENTIFIC AND TECHNICALKhorunov V.F., Voronov V.V. and Maksymova S.V. Brazing oftitanium alloys by using aluminium-base filler alloys .............................. 2

Markashova L.I., Akhonin S.V., Grigorenko G.M., KruglenkoM.G., Kushnaryova A.S. and Petrichenko I.K. Structure andproperties of welded joints on titanium alloys containing siliconadditions ............................................................................................. 6

Borisov Yu.S., Khaskin V.Yu., Vojnarovich S.G., Kislitsa A.N.,Tunik A.Yu., Adeeva L.I., Kuzmich-Yanchuk E.K., BernatskyA.V. and Siora A.V. Combined laser-microplasma claddingwith powders of Ni—Cr—B—Si system alloys ........................................... 16

Zhdanov L.A., Duchenko A.N., Goncharov I.A., Galinich V.I.,Zalevsky A.V. and Osipov N.Ya. Thermodynamic analysis ofslag melts in manufacture of fused welding fluxes ............................... 23

INDUSTRIALLobanov L.M., Pashchin N.A., Cherkashin A.V., Tkachuk G.I.,Savitsky V.V., Mikhoduj O.L., Shiyan K.V., Levchuk V.K.,Zhyginas V.V. and Lyashchenko A.P. Repair welding ofintermediate cases of aircraft engines from high-temperaturemagnesium alloy ML10 with application of electrodynamictreatment ........................................................................................... 28

Belokon V.M. and Koroteev A.O. Procedure for calculation ofdimensions of nozzles in welding using two separate gas jets .............. 33

Pismenny A.S., Pentegov I.V., Kislitsyn V.M., Stemkovsky E.P.and Shejkovsky D.A. Braze-welding with weld metal peeningduring its solidification ........................................................................ 37

Yakushin B.F. Comparative analysis of ISO 18841:2005standard and RF 26389—84 standard on evaluation to hotcrack resistance in welding ................................................................. 41

Dolinenko V.V., Skuba T.G., Vashchenko O.Yu. and LutsenkoN.F. Multichannel microprocessor controller for data collectionfrom thermocouples ............................................................................ 45

Makovetskaya O.K. Technological innovations – basis forincrease of competitiveness of the U.S. welding production ................. 48

NEWSIvanova O.N., Zelnichenko A.T., Kunkin D.D., Perekrest V.V.and Todorenko V.A. Experience of application of HF electricwelding apparatus EK-300M1 in surgery ............................................. 53

Scientific-Technical Conference «Modern Problems ofMetallurgy Technology of Welding and Surfacing of Steels andNon-Ferrous Metals» .......................................................................... 56

© PWI, International Association «Welding», 2012

English translation of the monthly «Avtomaticheskaya Svarka» (Automatic Welding) journal published in Russian since 1948

International Scientific-Technical and Production Journal

Founders: E.O. Paton Electric Welding Institute of the NAS of Ukraine Publisher: International Association «Welding» International Association «Welding»

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November2012# 11

BRAZING OF TITANIUM ALLOYSBY USING ALUMINIUM-BASE FILLER ALLOYS

V.F. KHORUNOV, V.V. VORONOV and S.V. MAKSYMOVAE.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

Investigations on brazing of titanium alloy samples by using different compositions of aluminium filleralloys were carried out. Silicon-free aluminium filler alloys were found to be acceptable for producing thebrazed joints on titanium alloys. The 670—690 °C brazing temperature range is optimal for the selectedfiller alloys.

Keywo r d s : brazing, titanium alloys, aluminium al-loys, commercial brazing filler alloys, wetting, micro-structure, mechanical properties

Since the 1960s, Al-base filler alloys have beenwidely used for brazing of titanium alloys. Purealuminium or alloys of the Al—Si, Al—Si—Cu andAl—Mg systems are mainly applied as brazingfiller alloys [7]. Compositions of some Al-basefiller alloys are given in Table 1.

Key advantages of aluminium filler alloys arelow melting temperature, low specific weight,good compatibility with titanium alloys basemetal and, in particular, good wetting and flow-ing into the gap. Therefore, special considerationhas been given to the aluminium filler alloyssince the time when the Ti-base alloys have foundapplication in aerospace engineering.

An important drawback of the Al-base filleralloys is their active reaction with the base metal.Even a relatively short time of contact of titaniumwith molten aluminium may lead to a deep ero-sion of the base metal. Silicon is added to theAl-base filler alloys to reduce reactivity of pure

aluminium and decrease the brazing temperature(hence, decrease the probability of formation ofintermetallics). But in this case silicides mayform at the titanium alloy—filler alloy interface.However, the main problem is the Al2O3 film onthe aluminium filler alloys, which prevents theirspreading over the base metal.

Despite a large amount of the investigationsconducted in Eastern Europe and particularly inUkraine to study brazing of titanium by usingaluminium filler alloys, brazing of titanium withthis type of the filler alloys failed to receive ac-ceptance. There are publications on developmentof new aluminium filler alloys for brazing of ti-tanium alloys [7], this evidencing the industrialdemand for commercial medium-melting pointfiller alloys for brazing of titanium and its alloys.

Wide application of aluminium filler alloysin this case is hindered by a low strength of theresulting brazed joints, which is much lower thanthat of the joints brazed with titanium filler al-loys. One of the promising areas of using alu-minium filler alloys is brazing of lamellar-ribbedthin-walled structures and thin-walled honey-

© V.F. KHORUNOV, V.V. VORONOV and S.V. MAKSYMOVA, 2012

Table 1. List of standard aluminium-base brazing filler alloys

Grade of filler alloyManufacturing

countryComposition of filler alloy Tbr, °С

AD1 USSR Al—0.4Si—0.3Fe 665

AL2 Same Al—13Si 560—700

AVCON 48 USA Al—4.8Si—3.8Cu—0.2Fe—0.2Ni 610—680

AA3003 Same Al—1Mn—0.6Si—0.7Fe 660—670

TiBrazeAl-600 » Al—12Si—0.8Fe 590—610

TiBrazeAl-630 » Al—1.5Mg—4Cu—2Ni 630—660

TiBrazeAl-640 » Al—(4.4—5.2)Mg—(0.7-1)Mn—0.2Cr 640—660

TiBrazeAl-642 » Al—5.3Si—0.8Fe—0.3Cu—0.2Ti 650—680

TiBrazeAl-645 » Al—(4.3—5.5)Mg—0.25Si—0.4Fe—0.2Ti—0.2Cr 640—660

TiBrazeAl-655 » Al—6.3Cu—0.3Mn—0.2Si—0.2Ti—0.2Zr 650—670

TiBrazeAl-665 » Al—2.5Mg—0.2Si—0.4Fe—0.2Cr 660—680

2 11/2012

comb panels for aerospace engineering, where arelatively low strength of the seams is acceptable,which is confirmed, in particular, in study [7].The choice of the aluminium filler alloys is fa-voured by their good wetting and spreading overthe titanium substrate at a comparatively lowtemperature, as well as the possibility of achiev-ing a low level of erosion of the base metal inbrazing. Therefore, all drawbacks in this case aresurpassed by the advantages, such as a lower cost,higher affordability and better workability of thealuminium filler alloys compared to the titaniumand silver ones.

The purpose of this study was to generate dataon advantages and drawbacks of different com-positions of aluminium filler alloys for brazingof titanium, as well to compare modern commer-cial and experimental filler alloys with widelyused aluminium alloys AD1 and AMg6.

Low titanium alloy OT4 was used as a basemetal. Two groups of aluminium filler alloys forbrazing of titanium were investigated: in the firstgroup silicon was used as a depressant, and inthe second group the depressant was magnesium.

The first group included standard alloy AL2,modern commercial filler alloy TiBrazeAl-642and experimental alloys Al—12Si—1Mg, Al—12Si—0.3Li and Al—5Si—1.5V produced by the powdermetallurgy method. The second group includedalloy AMg6 and modern commercial filler alloyTiBrazeAl-665. Low alloy AD1 was investigatedfor comparison.

Experiments on selection of optimal parame-ters of heating for brazing were carried out invacuum furnace SGV 2,4-2/15-I3 at a vacuumlevel of 5⋅10—5 mm Hg. For additional cleaningof the brazing atmosphere, brazing was per-formed in vacuum in the titanium container witha getter.

Table 2 gives contact angles of the alloys ona substrate of titanium alloy OT4, which weremeasured by using software AutoCad 2002LT.Increase in the temperature of brazing of titaniumalloys was accompanied by substantial improve-ment of wetting and spreading of the filler alloysover the substrate. However, it should be notedthat Si-containing filler alloys TiBrazeAl-642,Al—13Si and Al—12Si—0.3Li featured a poorspreading over the surface of the titanium sam-ples up to a temperature of 700 °C (at 740 °C,spreading of all the filler alloys was so high thatit caused flowing out of a filler alloy from thegaps, the contact angle in this case being approxi-mately 0°). At the same time, filler alloy AD1and the Mg-containing filler alloys (AMg6 andTiBrazeAl-665) satisfactorily wetted titaniumeven at 670 °C.

Metallographic examinations of the brazedjoints made by using the Mg-containing alu-minium filler alloys showed the presence of a

Figure 1. Microstructure of fillet region of the brazed jointon titanium alloy OT4 made by using filler alloy AMg6

Table 2. Dependence of contact angles on brazing temperature

Filler alloy TL, °С

Temperature of heating for brazing, °C

600 630 670 700

Contact angles, deg

AD1 (Al—0.4Si—0.3Fe) 660 — — 60 ~15

АМg6(Al—6Mg—0.6Mn—0.4Si—0.4Fe—0.1Ti)

632 — — 20 7—10

TiBrazeAl-642(Al—5.3Si—0.8Fe—0.3Cu—0.2Ti)

630 — — 40 ~10

AL2 (Al—13Si) 578 90 90 55 ~25

TiBrazeAl-665(Al—2.5Mg—0.2Si—0.4Fe—0.2Cr)

650 — — 25 8—10

Al—12Si—0.3Li 580 90 60 45 ~10

Al—12Si—1Mg 575 — 85 40 ~15

Al—5Si—1.5V 630 — — 40 ~10

11/2012 3

continuous intermetallic interlayer at the filleralloy—base metal interface (Tbr = 685 °C, vac-uum – 5⋅10—5 mm Hg, t = 3 min). Compositionof the interlayer varied from (wt.%) 48.67Al—47.95Ti—1.05Si—0.57Mn in the fillet region (seespectrum 1 in Figure 1; Table 3) to 72.68Al—20.75Ti—1.33Mg—0.74Si—0.36Mn (spectrum 2 inFigure 2; Table 4). In the first case it corre-sponded approximately to a composition of in-termetallic compound TiAl2, and in the secondcase – to TiAl3.

Also, one should note a low content of mag-nesium in the brazed seams, i.e. maximum1.5 wt.% (see Tables 3 and 4). This can be ex-plained by evaporation of magnesium from theseam metal during heating and melting of a filleralloy in vacuum. Very likely that it is this factthat caused destruction of the aluminium oxidefilm on the surface of the filler alloy, which madewetting of the base metal with the filler alloymelt much easier. The destructed oxide film wasdistributed over the entire seam (see oxygen con-tent in spectra 1—9 in Figure 1 and Table 3;

spectra 5—9 in Figure 2 and Table 4), except forthe intermetallic interlayer at the filler alloy—base metal interface.

Light phase inclusions along the seam axiswere compounds of aluminium with iron and sili-con, which were present in alloy AMg6 in smallquantities.

Metallographic examinations of the jointsbrazed by using the Si-containing aluminiumfiller alloys showed that the brazed joints werecharacterised by a poor quality and the presenceof cracks in the seams and fillet regions. Solidi-fication of silicide in the form of a continuousstrip was observed along the seam on both inter-faces with the base metal. Such peculiarities offormation of the brazed seams did not allow pro-viding of sound brazed joints and avoiding for-mation of silicides and cracks (Figure 3).

Strength tests of the overlap joints on alloyOT4 brazed by using commercial filler alloys Ti-BrazeAl-665 and TiBrazeAl-642, as well as alloysAD1 and AMg6 were carried out to evaluate thelevel of strength of the brazed joints. Thicknessof the filler alloy foils was 100 μm for TiBrazeAl-665 and TiBrazeAl-642, and 60 μm for AD1 and

Figure 2. Microstructure of region of the brazed joint ontitanium alloy OT4 made by using filler alloy AMg6

Figure 3. Microstructure of region of the brazed joint ontitanium alloy OT4 made by using filler alloy Al—5Si—1.5V

Table 3. Chemical heterogeneity of fillet region of the brazedjoint on titanium alloy OT4 made by using filler alloy AMg6,wt.%

Spectrum number

O Mg Al Si Ti Mn Fe

1 1.76 — 48.67 1.05 47.95 0.57 —

2 1.62 — 50.89 1.06 45.83 0.60 —

3 1.13 0.43 97.01 — 0.79 0.64 —

4 0.77 0.49 97.84 — 0.25 0.65 —

5 1.24 0.28 97.78 — 0.13 0.57 —

6 3.01 0.15 95.68 — 0.15 0.69 0.32

7 1.07 0.53 97.01 — 0.77 0.62 —

8 1.78 0.71 96.59 — 0.28 0.63 —

9 1.31 0.71 96.71 — 0.64 0.63 —

Table 4. Chemical heterogeneity of region of the brazed joint ontitanium alloy OT4 made by using filler alloy AMg6, wt.%

Spectrumnumber

O Mg Al Si Ti Mn Fe

1 — 0.46 64.12 1.32 33.55 0.30 0.25

2 4.14 1.33 72.68 0.74 20.75 0.36 —

3 — — 54.42 1.34 43.83 0.41 —

4 — 0.75 68.94 1.10 28.78 0.43 —

5 1.16 1.43 96.22 — 0.60 0.59 —

6 1.12 1.25 96.35 — 0.74 0.54 —

7 1.23 1.46 96.47 — 0.30 0.54 —

8 1.64 0.96 81.16 0.48 0.30 1.99 10.13

9 1.15 1.05 85.72 0.46 0.30 1.66 9.66

10 — — 4.49 — 94.95 0.56 —

11 — — 3.59 — 95.90 0.51 —

4 11/2012

AMg6. A filler alloy in the form of a foil wasplaced in the gap between the samples brazed.The time of holding at a brazing temperature was3 min, and the brazing temperature was 685 °C.Additionally, brazing of the samples by using theAMg6 filler alloy was carried out at a tempera-ture of 720 °C. The mechanical test results areshown in Figure 4.

Based on the data presented, it can be notedthat strength of the joints made by using Mg-containing filler alloys TiBrazeAl-642 and AMg6was almost identical and equal to 82—83 MPa,whereas strength of the joints made by usingSi-containing filler alloy TiBrazeAl-665 was low,which could be due to solidification of silicidein the form of a continuous strip at the filleralloy—base metal interface.

It should be noted that the evaluated strengthvalue (83 MPa) of the joints brazed by usingcommercial filler alloys TiBrazeAl-665 (Al—2.5Mg—0.3Cr) turned out to be lower than thatclaimed by the manufacturer (about 98 MPa)[7]. The attempts to achieve the claimed valuesfailed, and after changing the configuration ofthe samples brazed, which was aimed at decreas-ing the bending component of stresses in sheartests, the determined strength value of the brazedjoints was the same 83 MPa. Probably, in ourexperiments we omitted some know-how of theauthors.

Increase in the brazing temperature had anextremely negative effect on strength of the jointsbrazed with aluminium filler alloys. For example,a twofold decrease in strength was revealed inbrazing with alloy AMg6 at a temperature of720 °C (see pos. 3 in Figure 4). In this case thedecrease can be explained by growth of the Ti3Alinterlayer because of intensification of the reac-tivity of aluminium with respect to titanium withincrease in the temperature and extension of thetime of contact of the molten filler alloy withthe titanium substrate.

The obtained strength value (about 83 MPa)of the brazed joints made by using aluminiumfiller alloys is sufficient for brazing of lamellar-ribbed structures and sheet parts with a largecontact area. The main advantage of the alu-minium filler alloys in this case will be, as men-tioned above, the workability, low cost and af-fordability.

Analysis of the results obtained shows thatthe 680—690 °C brazing temperature is acceptablefor producing the brazed joints on titanium alloysby using the Si-free aluminium filler alloys, suchas alloys AD1, AMg6 and TiBrazeAl-642. Thetime of holding in brazing of titanium by usingthe above filler alloys should be as short as pos-

sible to prevent formation of brittle intermetallicinterlayers.

CONCLUSIONS

1. Si-free brazing filler alloys, e.g. AD1, AMg6,TiBrazeAl-642, were found to be acceptable forproducing the brazed joints on titanium alloys. Thebest results were obtained with the Al—Mg systembased filler alloys (AMg6, TiBrazeAl-642).

2. The 680—700 °C brazing temperature rangeis optimal for the chosen filler materials. The hold-ing time in brazing of titanium with the given filleralloys should be as short as possible to preventformation of brittle intermetallic interlayers.

3. When using filler alloys based on the Al—Sisystem, formation of silicides, in addition to theAl- and Ti-base intermetallic interlayers, occursin the brazed seams. They have the form of acontinuous strip propagating along the seam onthe side of the base metal, this leading to origi-nation of defects in the form of cracks.

1. Lashko, N.F., Lashko, S.V. (1977) Brazing of met-als. Moscow: Mashinostroenie.

2. Wells, R.R. (1975) Low temperature large-area braz-ing of damage tolerant structure. Welding J.,54(10), 348—356.

3. Kimbal, C.E. (1980) Aluminum brazed titaniumacoustic structures. Ibid., 59(10), 26—30.

4. Nesterov, A.F., Studenov, G.V. (1987) Brazing ofoverlap titanium joints with aluminium filler alloys.In: Proc. of Seminar on Increase in Quality and Ef-ficiency of Welding Production at Moscow Enter-prises (Moscow, 1987).

5. Nesterov, A.F., Dolgov, Yu.S., Telkov, A.M. (1985)Brazing of titanium structures with aluminium filleralloys. In: Filler alloys for brazing of current materi-als. Ed. by A.A. Rossoshinsky. Kiev: PWI.

6. Sokolova, N.M., Perevezentsev, V.N. (1989) Brazingwith aluminium filler alloys: Progress. In: Advanceedbrazing methods: Transact. Ed. by V.F. Khorunov.Kiev: PWI.

7. Shapiro, A.E., Flom, Y.A. Brazing of titanium attemperatures below 800 °C: Review and prospectiveapplications. http://www.titanium-brazing.com/publications/DVS-Manuscript_1020-Copy-19-07.pdf

Figure 4. Strength of the overlap joints on alloy OT4 (hold-ing time – 3 min) made by using the following filler alloys:1 – TiBrazeAl-642; 2, 3 – AMg6; 4 – AD1; 5 – Ti-BrazeAl-665 (1, 2, 4, 5 – Tbr = 685; 3 – 720 °C)

11/2012 5

STRUCTURE AND PROPERTIES OF WELDED JOINTS ONTITANIUM ALLOYS CONTAINING SILICON ADDITIONS

L.I. MARKASHOVA, S.V. AKHONIN, G.M. GRIGORENKO, M.G. KRUGLENKO,A.S. KUSHNARYOVA and I.K. PETRICHENKO

E.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

Structural-phase transformations in specimens of electron beam welded joints on two experimental heat-resistant pseudo- α and α + β multi-component titanium alloys containing silicon additions were investigated.Specific (differentiated) contributions of different types of structures and phase formations in the near-weldzone to strength values and distribution of local internal stresses in the welding zones under investigationwere analytically estimated.

Keywo r d s : heat-resistant titanium alloy, structuralstate, phase formations, microdiffraction reflections,strength characteristics, local internal stresses

Compared to aluminium alloys, steels and nickelsuperalloys, high values of strength, specificstrength and corrosion resistance of titanium al-loys over a wide temperature range favour theirincreasingly wider application in aircraft andspace engineering, ship building, chemical indus-try, etc. The use of titanium alloys grows due tohigh reliability of this class of materials at in-creased and high (of the order of 600—650 °C)temperatures, as well as in high-temperature andaggressive environments, this allowing replace-ment of parts and components of steels and otherstructural materials by titanium ones (parts ofcases of rocket engines and nuclear power plants,disks and blades of compressors, steam turbines,turbine and gas-turbine engines, heat exchangers,etc.). Heat-resistant titanium alloys are receivingan increasing acceptance in motor car construc-tion, this leading to a substantial increase inpower of automobile engines [1—3].

However, complication of service conditionsrelated to increase in the level of working tem-peratures and necessity to extend the life of partsand mechanisms requires not only improvementof composition and technology of treatment ofinitial materials, but also finding a solution tothe problem of their weldability. The latter is ofspecial importance in manufacture of long andcomplex-configuration structures, as well as inrepair-and-renewal operations, including, for ex-ample, reconditioning of worn-out engine blades.

As increase in service properties and level ofworking temperatures of any structure can beachieved, first of all, by appropriate alloying, aswell as by providing the required structural state

of the employed metals, alloys and their weldedjoints, the focus in this study is on conductingmore comprehensive investigations of structural-phase changes depending on alloying with sili-con, and on evaluating the chemical composi-tion → structure → properties relationship fortitanium alloys and their welded joints.

In this connection, considering the complexityof the processes and mutual effect of alloyingand phase formation under different technologi-cal conditions of the thermal-deformation effect(welding, heat treatment), it seems expedientnot only to perform appropriate experimentalstudies of structural-phase changes (chemicalcomposition, character of grain, sub-grain anddislocation structure, and phase precipitates dif-fering in composition, morphology and distribu-tion) under certain welding conditions, but alsoto evaluate the effect of specific structural-phasecomponents on changes in mechanical charac-teristics of the welded joints that are most sig-nificant for service conditions, such as strength,ductility and crack resistance values. This willmake it possible to determine the role of struc-tural and phase components not only in strength-ening of metal, but also as a factor affecting theprocesses of accumulation of local internalstresses, value and extent of this type of stresses,as well as the possibility of their plastic relaxa-tion, which is an indicator of crack resistance ofa material under service conditions.

Materials and procedures. The investigationobjects in this study are electron beam welded(EBW) joints on two heat-resistant multi-com-ponent titanium alloys. Both alloys contain sili-con as an alloying element, and belong to pseudo-α (alloy 1) and α + β (alloy 2) titanium alloys(Table).

© L.I. MARKASHOVA, S.V. AKHONIN, G.M. GRIGORENKO, M.G. KRUGLENKO, A.S. KUSHNARYOVA and I.K. PETRICHENKO, 2012

6 11/2012

The basic experimental information on struc-tural-phase composition of metal of a weldedjoint was generated by using optical, analyticscanning microscopy (SEM-515, PHILIPS, Hol-land) and microdiffraction transmission electronmicroscopy (JEM-200 CX, JEOL, Japan) withaccelerating voltage of 200 kV. Thin foils fortransmission microscopy were prepared by thetwo-stage method, i.e. preliminary electropolish-ing and subsequent multiple ion thinning by ion-ised argon flows in a specially developed unit[4]. The latter allowed not only widening theinvestigation fields (increasing statistics), butalso making all structural and phase componentsof a material being analysed «transparent» forelectrons.

Investigation results. The optical metallogra-phy methods were used to reveal structure, pres-ence and arrangement of cold cracks in the EBwelded joints on two experimental titanium al-loys in the most problematic zone of a weldedjoint, i.e. HAZ [5—7], which in fact is a near-weldzone (NWZ), whose size and structure are deter-mined by the thermal cycle of welding, and wherethe most dramatic changes in structure are ex-pected to take place, allowing for high coolingrates characteristic of EBW.

Cold cracks were found to form in the weldedjoints on the investigated alloys after welding,the rate of formation of this type of cracks in thewelded joints on experimental alloy 1 being muchhigher (Figure 1, a, b) than in the welded jointson alloy 2.

Also, as shown by metallographic examina-tions of structure, coarse equiaxed polyhedral

primary β-grains up to 0.5 mm in size form inNWZ of the welded joints on alloy 1 (Figure 1,c). In NWZ of the welded joints on alloy 2 theprimary structure is heterogeneous: along withlarge regions of polyhedral grains 0.2 mm in size(Figure 1, d), there are regions of fine 20—60 μmequiaxed grains surrounded by coarse grains(Figure 1, e). Formation of chains of fineequiaxed grains was observed also in the HAZregions located at a distance from the weld (Fi-gure 1, f). As a rule, they extend along the basemetal rolling direction (normal to the weld axis).Often, location of fine grains coincides with lo-calisation of clusters of dispersed precipitates,most probably silicide ones. Intragranular struc-ture in NWZ of alloy 1 consists of a coarse-acicu-lar α′-phase. In NWZ of alloy 2, the martensiticα′-phase has a fine-acicular structure (see Figu-re 1, c, d). In addition to the martensitic phase,NWZ of both alloys may contain the retainedβ-phase, the amount of which, according to thechemical composition, is very insignificant in al-loy 1, and higher in alloy 2 than in alloy 1.

More detailed structural-phase examinationsof HAZ of the welded joints on titanium alloysby using microdiffraction transmission electron

Figure 1. Microstructures of HAZ metal on experimental heat-resistant alloys after EBW: a, b – alloy 1, cracks in HAZmetal; c – alloy 1, NWZ; d, e – alloy 2, NWZ; f – alloy 2, HAZ region located at distance from the weld

Chemical composition of experimental heat-resistant alloys

AlloyContent of alloying elements, wt.% Coefficient

of stabilityof β-phase KβAl Sn Zr Mo V Nb Si

1 5.2 3.3 4.2 0.1 0.6 0.8 0.6 0.07

2 4.3 4.4 6.0 1.6 0.7 4.3 0.4 0.33

11/2012 7

microscopy were carried out to determine com-position of the forming phases, as well as theirsizes, morphology and structural zones of theirlocalisation (internal volumes or grain-boundaryregions).

Welded joint on experimental heat-resistantalloy 1. Structure of NWZ of the EB weldedjoint on alloy 1 consists mainly of the laminatedα′-phase and a very small amount of the lami-nated β-phase, which differ in length ll of a sub-micron sized (approximately from 0.3 to 1.5 μm)form with cross section hl (Figure 2, a). More-over, the laminated structural components differgreatly in their internal structure. The major partof this type of structures (consisting mostly ofthe α′-phase, according to the microdiffractionanalysis) is characterised by a minimal disloca-tion density (ρ ~ 109 cm—2) in the internal volumeof the uniformly distributed laminae. The otherpart of the laminated structures (their quantitybeing much lower) radically differs both in dis-location density and distribution. For example,in this type of the laminated structures the dis-location density is higher approximately by anorder of magnitude (ρ ~ (7—8)⋅1010 cm—2). Thedistribution of crystalline lattice defects in somecases is more or less uniform (Figure 2, a, b),whereas in other cases the complex dislocationconfigurations in the form of blocks or cells, as

well as an intralaminar dispersed (ds ~ 0.1 μm)sub-structure (Figure 3, a) are detected. Thestructure with a clearly defined intralaminar sub-structure is most pronounced in the dark fieldimaging mode (Figure 3, c).

It should be noted that structures with a highdislocation and phase precipitate density corre-spond not only to the β-phases, but also partiallyto the α′-phases.

Examinations of thin foils allowed generatingthe detailed information on the phase precipitatesforming in the welded joint, which differ in size,morphology, stoichiometric composition and lo-calisation zones (along the boundaries of thelaminated structures, in internal volumes, in sub-structure, etc.).

Phase precipitates of fine sizes (dPh.P ~ 0.01—0.10 μm) forming in narrow grain-boundary in-terlayers and along the interlaminar boundaries(see Figure 2, c), the composition of which cor-responds mainly to stoichiometry Ti5Si3 (Figu-re 3, b), are most distinct. The fine phase pre-cipitates form also in internal volumes of the α′-and β-laminae, in the bulk of which fragmenta-tion of the intralaminar structure and formationof sub-structures take place (Figure 3, a—c). Thephases forming in this type of the structures aredetected primarily in the zones of intralaminarsub-structural boundaries, and are characterised

Figure 2. Microstructure of experimental alloy 1, NWZ: a – well-defined orientation of laminae of mainly α-componentof structure at comparatively low density and uniform distribution of dislocations (lamina width hl ~ 0.3—1.5 μm),×20,000; b, c – phase formation in internal volumes and in boundary regions of α-laminated structures, ×30,000

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by the finest sizes, i.e. dPh.P ~ 0.01—0.02 μm. Ascan be seen, such phases are the phases borderingthe sub-structure. In addition to the fine equiaxedphase precipitates, there are also precipitates ofan extended form, when lPh.P >> hPh.P at ll ~~ 0.7—0.8 μm, propagating along this type of thesub-structural boundaries (Figure 3, a, b).Stoichiometric composition of the fine phase pre-cipitates bordering the intralaminar sub-struc-ture becomes a bit wider: in addition to the notedTi5Si3 composition, there are also phases of othercompositions, including such elements as alu-minium and zirconium, i.e. Ti3Al and Ti2Zr3Si3phases (Figure 3, a—c; Figure 4, a).

The most active development of phase forma-tion is characteristic of the laminated structuresof comparatively coarse (in cross section) sizes(hl ~ 0.4—1.5 μm). Besides, the active phase for-mation in such zones is accompanied by occur-rence of the following important factors. Firstly,coarsening of the phase formations takes place,i.e. size of the phase formations dPh.P amountsto about 0.1—0.2 μm, this being an order of mag-nitude higher than size of the intralaminar sub-boundary phases observed in the laminated struc-tures of a smaller cross section (see Figure 4, a).Secondly, no ordering can be seen in distributionof coarse, mainly silicide phases in the bulk ofthe massive α′-laminae: the forming phases are

distributed chaotically, and they are not relatedeither to structural boundaries, or grain and sub-grain boundaries. Moreover, formation of intra-volume phases in the said cases is accompaniedby a substantial increase, i.e. up to (7—8)⋅1010 cm—2, of the dislocation density in thephase formation zone propagating along the en-tire length of the laminae (Figure 4, b, c). There-fore, a distinctive feature of the structure of themetal under investigation is formation of ex-tended, special α-lamina structural zones satu-rated with coarse globular phase precipitates sur-rounded by dense dislocation clusters.

As follows from the results of investigationsof the dislocation structure and phase formationprocesses, a substantial difference between thestructural-phase states of the α′- and β-laminatedstructures is observed in the welded joints onexperimental alloy 1. There occurs parallel for-mation of the laminated structures dramaticallydiffering in their structural-phase states, such asalmost dislocation-free laminae containing nophase precipitates, along with laminae charac-terised by a high dislocation density and satura-tion of the internal volumes with chaotically dis-tributed precipitates of a rather coarse size. It islikely that formation of the substantially graded(as to phase precipitates and dislocation density)laminae is attributable to the type of the crys-

Figure 3. Microstructure of experimental alloy 1, NWZ: a – fine structure of laminae with sub-structure, ×37,000; b –microdiffraction reflection; c – dark-field image of specific (marked with arrows in Figure 3, a) phase formations, ×3000

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talline lattice corresponding to the β- and α-for-mations in titanium alloys. For instance, the β-phase having the bcc lattice (comprising up to48 sliding systems) has an almost unlimited pos-sibility for initiation, sliding and redistributionof dislocations, which are known to serve as ac-tive channels for transportations of alloying ele-ments and, hence, activation of the phase forma-tion processes. The α-structure having the hcplattice is characterised by a very limited quantityof the sliding systems. Predominantly, this is onebasal (0001) plane, and deformation in metalwith this type of the lattice is realised due totwinning, which hampers dislocation initiationand sliding and, therefore, phase formation.

Most probably, it is different peculiarities inrealisation of the deformation processes (throughdislocation sliding or twinning) and, as a result,different phase formation possibilities for themain phase components (α- and β-phases) thatexplain formation of the extended laminatedstructures characterised by sharp gradients of thedislocation density and saturation with phaseprecipitates. The presence of the graded struc-tural-phase formations, which are substantiallydifferent in the quantity and degree of dispersionof the silicide phases, including in dislocationdensity, is likely to serve as a base for formationin metal of this type of the corresponding sharplygraded mechanical characteristics, such as gradi-ents of strength properties (σ0.2 and σt) in therelated laminated structures.

Therefore, it was found that NWZ of alloy 1is characterised by the presence of the extendedα′- and β-laminated phase formations, sharplygraded in dislocation density, as well as in quan-tity and size of the forming silicide and intermet-allic phase precipitates:

• α′ – laminated phase components (hcp lat-tice) characterised by a minimal intralaminar dis-location density and an insignificant quantity ofphase precipitates in laminae;

• β – laminated structures (bcc lattice) anda small part of the α′-phase characterised by adramatic increase in the general dislocation den-sity, formation of the sub-structure, very inten-sive development of the phase formation proc-esses (growth of size and quantity of phases) anddistribution of the silicide and intermetallicphases in zones of the dislocation clusters.

Welded joint on experimental heat-resistantalloy 2. Metal structure in NWZ of the EBwelded joint on experimental alloy 2, similar toexperimental alloy 1, is represented by differentphases (α′- and β-phases), which differ both insize and fine structure of the phase formationsand in size and distribution of the silicide andintermetallic precipitates originating during thewelding process.

For example, cross section size hl of laminaeof the martensitic α′-phase is much smaller (ap-proximately 2—3 times) compared to that of thelaminated structures of the corresponding zoneof the welded joint on experimental alloy 1, and

Figure 4. Microstructure of experimental alloy 1, NWZ: a – phase formation in β-laminae, ×50,000; b – extendeddramatic gradients of dislocation density along the laminated structures, ×30,000; c – combined microdiffraction reflec-tions of specific phases in α′-lamina structures

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is equal to 0.2—0.5 μm (Figure 5, a). In addition,no dramatic changes in thickness of the laminaeare observed. In this case, and this should beemphasised, structure of the α′- and β-phases ischaracterised by the presence of the acicular andfine intralaminar sub-structure. The dislocationdensity equal to ρ ~ (8—9)⋅1010 cm—2 is uniformlydistributed.

As to the phase precipitates, structural exami-nations and parallel analysis of microdiffractionreflections (Figure 5, b, d; Figure 6, b, d) showformation of primarily fine (0.01—0.02 × 0.02—0.06 μm) and comparatively more uniformly dis-tributed silicide and intermetallic phases in NWZof the welded joint on alloy 2, compared to thewelded joint on alloy 1. Moreover, the formingphases are distributed mainly in internal volumesof the laminated structures, first of all along thesub-structural boundaries, i.e. they are phase pre-cipitates that border the intralaminar sub-struc-tural elements (Figures 5 and 6). This characterof distribution of the fine phase precipitatesshould promote not only fixation of the formedintralaminar sub-structure, but also consolida-tion of the thus fixed structure up to a tempera-ture of dissolution of the grain-boundary distrib-uted phases. Besides, this type of the structuralstate (fine fragments with grain-boundary fixingphases) is more or less uniformly distributed inthe entire volume of the NWZ metal.

Analysis of microdiffraction reflections of thestructures being investigated reveals diversity ofstoichiometric compositions of the phase precipi-tates forming in NWZ of the joints on alloy 2.These are mostly phases of the Ti5Si3, Ti2Zr3Si3and Ti3Al types (Figure 5, d; Figure 6, b). Asseen, compositions of the precipitated silicidesand intermetallics hardly differ from those de-tected in NWZ of the welded joint on alloy 1.However, morphology of this type of the phases,their size and distribution are substantially dif-ferent. In the welded joint on alloy 2, silicidesand intermetallics are finer, have a rod-like orglobular shape (see dark-field image in Figure 6,b), and are distributed more uniformly in thebulk of metal, which seems to be caused by astructural state of the NWZ metal of the weldedjoint on this alloy, i.e. by a comparatively moreuniform and finer structure of α′-martensite.However, despite a more favourable change instructural-phase state of the NWZ metal on alloy2, including dispersion and uniformity of struc-ture, formation of fine precipitates along thestructural boundaries and absence of the lami-nated structure that is dramatically graded in itsstructural-phase state, the presence of a pro-nounced extension of the laminated structureswill lead, though to a smaller degree (comparedto the NWZ state in alloy 1), to decrease in

Figure 5. Microstructure of experimental alloy 2, NWZ: a, c – fine structure of laminated phases of the martensitictype (a – ×50,000; c – ×37,000); b, d – microdiffraction reflections of phase precipitates

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ductility values and, accordingly, to increase insusceptibility of the welded joint to cracking.

Therefore, NWZ of the welded joint on ex-perimental alloy 2 is characterised by formationof the extended phases of the laminated type(α′-martensite and β-phase) having, like in alloy1, a laminated morphology, but considerably dif-fering (approximately 2—3 times) in width of thelaminated structures, finer acicular α′-martensi-tic structure and intralaminar sub-structure, as

well as more uniform distribution of dislocationsin the entire volume of the NWZ metal.

Differences are observed also in the processof formation of the silicide and intermetallicphases: at a similar stoichiometric composition(like in case of alloy 1) the phases are smallerin size and are uniformly distributed in the entirevolume, their localisation occurring mainly alongthe sub-structure boundaries.

Additional fractographic examinations offractures of the EB welded joints on experimental

Figure 7. Microstructures of fracture surfaces on titanium alloys (×4020): a – brittle cleavage in laminated structureswith intravolume phase precipitates (welded joint on experimental alloy 1); b – quasi-brittle fracture in martensiticcomponent (welded joint on experimental alloy 2)

Figure 6. Microstructure of experimental alloy 2, NWZ: a, c – distribution of phase precipitates differing in morphologyand size, ×50,000; b, d – microdiffraction reflections of specific zones of phase precipitates

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alloys 1 and 2 showed that the fracture zone ofthe welded joints on experimental alloy 1 is char-acterised by the presence of regions of the ex-tended transcrystalline brittle cleavage in a di-rection of the laminated structures (Figure 7, a).In contrast to this, the welded joints on experi-mental alloy 2 feature a more homogeneous quasi-brittle fracture of the intragranular type (Fi-gure 7, b) with dispersed fragments (df ~ 2—5 μm) corresponding in size to sub-structuralcomponents in α′-martensite.

A substantial effect on quality of the weldedjoint is exerted by distribution and localisationof internal stresses in the HAZ metal of the alloysinvestigated. Stresses of this type related to non-uniformity of heating and structural-phase trans-formations lead to a dramatic decrease in ductil-ity, and in some cases to cold cracking, whichoccurs under the conditions of EBW of experi-mental heat-resistant alloys. Therefore, analysisof the role of different structural factors inducing

or blocking formation of internal stresses is alsoof an important practical interest.

The package of the conducted experimentalstudies made it possible, firstly, to analyticallyestimate specific (differentiated) contributionsof different structural-phase factors and parame-ters forming in welded joints of the investigatedalloys to changes in strength characteristics σ0.2,and, secondly, to reveal the structural factorsdetermining the character and distribution of in-ternal stresses τin, which are potential sources ofinitiation and propagation of cracks in the inves-tigated structural microregions [8—12].

Analytical estimates of strength σ0.2 weremade according to the Archard equation that in-cludes the known Hall—Petch, Orowan and otherdependences [13—20]:

ΣΔσ0.2 = Δσ0 + Δσs.s + Δσg + Δσs + Δσd + Δσd.s,

where Δσ0 is the resistance of the metal latticeto movement of free dislocations (friction stress

Figure 8. Contribution of different components of structural strengthening (grain, sub-grain, dislocation and dispersion):a – alloy 1; b – alloy 2; c – calculated value of yield strength Σσ0.2

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of the lattice or Peierls—Nabarro stress); Δσs.s isthe strengthening of solid solution with alloyingelements and impurities (solid solution strength-ening); Δσg, Δσs is the strengthening due to achange in size of grain and sub-grain (Hall—Petchdependences, grain and sub-grain strengthen-ing); Δσd is the dislocation strengthening causedby the inter-dislocation interaction; and Δσd.s isthe strengthening provided by the dispersed par-ticles according to the Orowan dependence (dis-persion strengthening).

It was shown as a result that the HAZ metalof the welded joint on experimental alloy 1 fea-tures the dramatically graded (approximately 1.8times) change in yield strength (Δσ0.2 ~ 570—1010 MPa) that depends on the structural-phasestate of the laminated structures. A dramatic in-crease in Δσ0.2, which is characteristic of the lami-nated structures with a high dislocation density(ρ ~ (7—8)⋅1010 cm—2) and most saturated withthe phase precipitates, leads to a growth of dis-location (Δσd ~ 250 MPa) and dispersion (Δσd.s ~~ 375—500 MPa) strengthening (Figure 8, a, c).

NWZ of alloy 2 is characterised by a highlevel and more uniform distribution of strengthproperties (Δσ0.2 ~ 910—1040 MPa) in the form-ing martensitic phases of the laminated type (Fi-gure 8, b, c), this being related to their finerstructure. In this case, a certain increase instrengthening is caused by dispersion of the sub-structure (Δσs ~ 530 MPa), and a comparativelyuniform increase in general dislocation densityin the bulk of metal leads to strengthening of anorder of Δσd ~ 360 MPa (Figure 8, b).

Furthermore, internal stresses τin in HAZ ofthe joints were determined by examinations ofthe dislocation structure [21, 22]:

τin = Gbhρ/[π(1 — ν)],

where G is the shear modulus; b is the Burgersvector; h = 2⋅10—5 cm is the foil thickness; ν isthe Poisson ratio; and ρ is the dislocation density.

The investigations conducted showed (Figu-re 9, a) that the HAZ metal of alloy 1 is char-acterised by a dramatically graded (approxi-mately 10 times) distribution of internal stresses,directed along the laminae (from 10—100 to 750—860 MPa), this being related to a change of thedislocation density in different types of the lami-nae, i.e. with low (ρ ~ 109—1010 cm—2) and high(ρ ~ (7—8)⋅1010 cm—2) dislocation densities. How-ever, there are also regions with an even higherlocal dislocation density (ρ ~ 2⋅1011 cm—2), wherelocal internal stresses τin/l amount to about2000 MPa.

HAZ of alloy 2 is characterised by a compara-tively uniform distribution of internal stresses

(τin ~ 860—970 MPa), this corresponding to auniform dislocation density (ρ ~ (8—9)⋅1010 cm—2)in the intralaminar structures (Figure 9, b).

However, both welded joints on alloy 1 andwelded joints on alloy 2 (though to a lower de-gree) feature a clear relationship of orientationof the distribution of internal stresses and thelaminated structures, which can be a cause offormation and propagation of cracks.

CONCLUSIONS

1. As established in the course of the comprehen-sive investigations of the welded joints on ex-perimental titanium alloys conducted at differentstructural levels (grain, sub-grain, dislocation),NWZ of the joints on alloys 1 and 2 is charac-terised by formation of the laminated-type ex-tended structures of the α′- and β-phase compo-nents with a similar morphology, but consider-ably differing in density and distribution of dis-locations, as well as in intensity of the processesof formation of phase precipitates of the silicideand intermetallic types.

2. In NWZ of the welded joint on pseudo- αalloy 1, the silicide phase formation occurs mostactively in few grains of the β-phase and in asmall part of the α′-laminae, which are charac-terised by a high dislocation density and forma-tion of a sub-structure. At the same time, the

Figure 9. Level of local internal stresses forming in lami-nated structures of NWZ of the welded joints: a – lami-nated structures graded in distribution of dislocation den-sity, and intravolume phase precipitates (experimental alloy1); b – martensitic laminated structures (experimental al-loy 2)

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major part of the α′-laminae is characterised bya low dislocation density, uniform distributionof dislocations and absence of silicides and inter-metallics in their bulk. Phase precipitates areobserved both in the grain-boundary interlayersand along the boundaries between the laminae.

3. The presence of the structural-phase forma-tions in NWZ of alloy 1, which are considerablydifferent in quantity and degree of dispersion ofthe silicide phases, and in dislocation density, isa base for formation of dramatically gradedstrength characteristics, as well as internalstresses in the adjoining laminated structures.

4. NWZ of the welded joints on (α + β) tita-nium alloy of the martensitic type is characterisedby formation of finer silicide and intermetallicphase precipitates in the α′- and β-phases, whichare mainly uniformly distributed in the bulk ofthe NWZ metal, i.e. along the sub-boundariesand boundaries of the fine martensitic α′-phase.

5. Analytical estimation of differentiated con-tribution of different structural-phase factors andparameters forming in the welded joints on theinvestigated alloys to changes in strength prop-erties (σ0.2) showed that a substantial change inyield strength σ0.2 of the adjoining laminatedstructures occurs in NWZ of the welded jointson alloy 1, i.e. from 570 MPa for the laminatedα′-phase with a low dislocation density to1010 MPa for the laminae with a high dislocationdensity and silicide precipitates. NWZ of alloy2 features a higher level and more uniform dis-tribution of strength properties (σ0.2 changesfrom 910 to 1040 MPa in the entire volume ofthe NWZ metal).

6. Estimation of changes in internal stressesτin in NWZ of the welded joints on the investi-gated alloys, made on a base of examinations ofthe dislocation structures, showed that distribu-tion of internal stresses in NWZ of the weldedjoint on alloy 1 is extremely non-uniform anddirected along the laminated structures (τinchanges from 10—100 to 750—860 MPa in thelaminae with a high and low dislocation densi-ties). Internal stresses in NWZ of the weldedjoint on alloy 2 are distributed more uniformly.However, fixation of the direction of the distri-bution of internal stresses and laminated struc-tures can serve as a cause of a directed propaga-tion of cracks.

7. To eliminate strength and internal stressgradients, it is necessary to achieve formation ofa homogeneous uniform dispersed structure.

1. Iliin, A.A., Kolachev, B.A., Polkin, I.S. (2009) Tita-nium alloys. Composition, structure, properties: Re-fer. book. Moscow: VILS—MATI.

2. Solonin, O.P., Glazunov, S.G. (1976) Refractory ti-tanium alloys. Moscow: Metallurgiya.

3. Chechulin, B.B., Ushkov, S.S., Razuvaeva, I.N. etal. (1977) Titanium alloys in machine-building. Mos-cow: Mashinostroenie.

4. Darovsky, Yu.F., Markashova, L.I., Abramov, N.P.et al. (1985) Procedure of thinning of dissimilarwelded joint samples for microscopic examinations.Avtomatich. Svarka, 12, 60.

5. Grabin, V.F. (1975) Principles of metals science andheat treatment of titanium alloy welded joints. Kiev:Naukova Dumka.

6. Moiseev, V.N., Kulikov, F.R., Kirillov, Yu.G. et al.(1979) Titanium alloy welded joints. Moscow: Me-tallurgiya.

7. Gurevich, S.M., Zamkov, V.N., Blashchuk, B.E. etal. (1986) Metallurgy and technology of welding oftitanium and its alloys. Kiev: Naukova Dumka.

8. Markashova, L.I., Grigorenko, G.M., Poznyakov,V.D. (2009) Influence of thermal cycles of weldingand external loading on structural-phase variationsand properties of joints of 17Kh2M steel. The PatonWelding J., 7, 18—25.

9. Markashova, L.I., Grigorenko, G.M., Arsenyuk,V.V. et al. (2002) Criterion of evaluation of me-chanical properties of dissimilar material joints. In:Proc. of Int. Conf. on Mathematical Modeling andInformation Technologies in Welding and RelatedProcesses (16—20 Sept. 2002, Katsiveli, Ukraine).Kiev: PWI, 107—113.

10. Markashova, L.I., Grigorenko, G.M., Poznyakov,V.D. et al. (2004) Structural approach to evaluationof mechanical properties in HAZ of steel and alloyjoints. In: Proc. of Int. Conf. on Mathematical Mod-eling and Information Technologies in Welding andRelated Processes (13—17 Sept. 2004, Katsiveli,Ukraine). Kiev: PWI, 174—179.

11. Markashova, L.I., Grigorenko, G.M., Poznyakov,V.D. et al. (2008) Structural factors determining theproperties of strength, plasticity and fracture ofwelded joints. In: Proc. of Int. Conf. on Mathemati-cal Modeling and Information Technologies inWelding and Related Processes (27—30 May 2008,Katsiveli, Ukraine). Kiev: PWI, 87—94.

12. Markashova, L.I., Grigorenko, G.M., Poznyakov,V.D. et al. (2009) Structural criterion for evaluationof strength, plasticity, crack resistance of metals, al-loys, composite materials and their welded joints. In:Proc. of 4th Int. Conf. on Fracture Mechanics ofMaterials and Strength of Structures (June 2009,Lviv, Ukraine). Lviv: FMI, 447—451.

13. Suzuki, H. (1967) About yield strength of polycrys-talline metals and alloys. In: Structure and mechani-cal properties of metals. Moscow: Metallurgiya.

14. Ashby, M.F. (1986) About Orowan stress. In: Phys-ics of strength and plasticity. Moscow: Metallurgiya.

15. Goldshtein, M.I., Litvinov, V.S., Bronfin, B.M.(1986) Physics of metals of high-strength alloys.Moscow: Metallurgiya.

16. Conrad, H. (1973) Model of deformation strengthen-ing for definition of grain size effect on metal flowstress. In: Ultrafine grain in metals. Ed. by L.K.Gordienko. Moscow: Metallurgiya.

17. Armstrong, R.V. (1973) Strength properties of met-als with ultrafine grain. Ibid.

18. Petch, N.J. (1953) The cleavage strength of polycrys-talline. J. Iron and Steel Inst., 173, 25—28.

19. Orowan, E. (1954) Dislocation in metals. New York:AIME.

20. Ashby, M.F. (1983) Mechanisms of deformation andfracture. Adv. Appl. Mech., 23, 117—177.

21. Koneva, N.A., Lychagin, D.V., Teplyakova, L.A. etal. (1986) Dislocation-disclination sub-structures andstrengthening. In: Theoretical and experimental in-vestigation of disclinations. Leningrad: LFTI.

22. Conrad, H. (1963) Effect of grain size on the loweryield and flow stress of iron and steel. Acta Metal-lurgica, 11, 75—77.

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COMBINED LASER-MICROPLASMA CLADDINGWITH POWDERS OF Ni—Cr—B—Si SYSTEM ALLOYS

Yu.S. BORISOV, V.Yu. KHASKIN, S.G. VOJNAROVICH, A.N. KISLITSA, A.Yu. TUNIK, L.I. ADEEVA,E.K. KUZMICH-YANCHUK, A.V. BERNATSKY and A.V. SIORA

E.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

Structural features of deposited layers produced by the combined laser-microplasma method using powdersof the Ni—Cr—B—Si system alloys were investigated. Technological advantages and drawbacks of a combi-nation of laser cladding and microplasma spraying were determined. It was shown that the developedcombined laser-microplasma method allows improving the quality of the deposited layers by preserving thekey advantages characteristic of the laser powder cladding process.

Keywo r d s : combined laser-microplasma cladding,self-fluxing nickel alloy, structure, phase composition,hardness, wear resistance

Different thermal spraying methods, such asflame, plasma and detonation ones, are appliedto deposit coatings of alloys of the Ni—Cr—B—Sisystem. Spraying and melting can be performedin one (gas powder cladding) or two successivestages (spraying with subsequent melting of thesprayed layer). The sprayed NiCrBSi coatingspreserve the main properties of the NiCrBSi alloy(wear and corrosion resistance), but lack the highadhesion strength (normally, less than 35—40 MPa). After melting the strength of adhesionof the NiCrBSi coating layer to the substrategrows to 70—75 MPa [1].

For a number of industrial problems it is de-sirable that the adhesion strength value be asclose as possible to strength of the base metal.In this connection, of an increasing interest nowis the process of laser melting of coatings. Ad-vantages of this process include thermal localityand minimal effect on the base metal, as well assmall (5—20 μm) size of the transition zone, whichminimises penetration of the base metal into thedeposited one and favours refining of structureof the material, this resulting in improvement ofmechanical properties. However, shrinkagecracks may form in laser melting as a result ofdramatically heterogeneous heating, especially ofcoatings more than 0.5 mm thick, as well as sub-sequent cooling [2, 3].

It is noted in studies [4—6] that drawbackscharacteristic of laser melting can be eliminatedby combining plasma and laser heating. One ofsuch processes, which integrates advantages oflaser cladding and microplasma spraying, is com-bined laser-microplasma cladding (CLMPC) [7].

It allows avoidance of drawbacks characteristicof laser cladding (formation of internal pores andmicrocracks), preparation of the workpiece sur-face directly during the process of deposition ofa material, and fusion of the deposited layerswith the base metal.

The purpose of this study was to investigatestructural peculiarities of the layers depositedwith powders of the Ni—Cr—B—Si system alloys(PG-12N-01 and PG-12N-02) by the CLMPCprocess, as well as to define technological advan-tages of combining the laser cladding and micro-plasma spraying processes.

Investigation procedure. Layers 0.3—1.2 mmthick were deposited on substrates of steels St3and 38KhN3MFA by the CLMPC method usingself-fluxing alloy powders (PG-12N-01 and PG-12N-02). Structure, phase composition and prop-erties of the layers were investigated. An inte-grated procedure comprising metallography (mi-croscope «Neophot-32» with digital photographyattachment), durometric analysis (LECO hard-ness meter M-400 with loads of 0.25, 0.5 and1 N) and X-ray phase analysis in monochromaticCuKα-radiation by using diffractometer DRON-UM1 was applied for investigation of the result-ing deposited layers. Graphite single crystalplaced on a path of the diffracted beam was usedas a monochromator. Diffraction patterns weremade by the step scan method in the 20° < 2θ << 90° angle range. The scan step was 0.05°, andthe time of exposure at a point was 3—7 s. Thedata of the diffractometry experiment were proc-essed by using software PowderCell 2.4 for full-profile analysis of X-ray spectra of a mixture ofpolycrystalline phase components.

Cracking index α the value of which was de-termined in percent was introduced to compare

© Yu.S. BORISOV, V.Yu. KHASKIN, S.G. VOJNAROVICH, A.N. KISLITSA, A.Yu. TUNIK, L.I. ADEEVA, E.K. KUZMICH-YANCHUK, A.V. BERNATSKY and A.V. SIORA, 2012

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the quantity of cracks in clad specimens. Theabsence of cracks in the deposited layer was takenas zero, and the network of cracks with a pitchof 1 mm was taken as 100 %. The investigationsshowed that this index can be estimated fromformula α = 3/L, where L is the distance betweenthe cracks, mm.

Investigation of wear resistance of the depos-ited layers was carried out by using friction ma-chine 2070 SMT-1 by the disk—pin method with-out lubrication. Before the investigation the40 mm diameter specimens were polished to sur-face finish Ra = 1.6 μm. Material of the matingbody was steel 45 heat treated to hardnessHRC 55. Prior to the tests, the surfaces weresubjected to running-in, the presence of whichwas fixed from stabilisation of the friction mo-ment in a pair. Relative wear resistance was de-termined from the loss of weight at sliding speedsof 1.3 m/s in a mode of stepwise loading, thetest time at each step being 15 s, and load being0.2 kN. Wear resistance of the clad specimenscompared to that of the base metal was deter-mined by using the in-house friction machine ac-cording to the cylinder—pin scheme by the dryfriction method. As this machine is a non-stand-ard development, the results obtained by usingit were regarded as relative. A specimen of steel38KhN3MFA after volumetric (furnace) harden-ing and heat treatment, having hardnessHRC 43—44, was chosen as the reference one.Wear of the specimens measured from a changein weight (in grams) was compared with that ofthe reference specimen. The time of friction wasincreased to improve accuracy of the measure-ment results. The mating body was made fromhard alloy T15 or steel 45 hardened to HRC 50—55. The specific pressure was set within 11—12 MPa, the rotation speed for a test specimenwas 50—1600 rpm, and the linear friction speedswere 0.4—15.0 m/s. In all the cases the specimenswere weighed before and after the friction testsby using analytical balance VPR-200 within0.0005 g.

Materials, equipment and principle of opera-tion of the devices. Self-fluxing nickel alloypowders PG-12N-01 and PG-12N-02 with a par-ticle size of 40—100 μm (TU 48-19-383—84), thechemical compositions of which is given in Ta-ble 1, were used as materials for cladding. Thepowders produced by atomisation in inert gashad particles with a regular round shape, closeto the spherical one. The fractional compositionof the powders was —40—+100 μm. Hardness of

the PG-12N-01 powder was HRC 36—45, andTmelt = 1080 °C. Hardness of the PG-12N-02 pow-der was HRC 46—55, and Tmelt = 1050 °C.

CO2-laser TR-100 (Rofin-Sinar, Germany)with a power of up to 10 kW was used as a laserradiation source. Radiation of this laser with apower of 2, 3 and 4 kW, combined with themicroplasma jet with a power of up to 1.5 kW,was used in the experiments.

The MPN-004 system with the MP-04 micro-plasmatron developed by the E.O. Paton ElectricWelding Institute of the NAS of Ukraine wasused to form the microplasma jet (design of themicroplasmatron is covered by the Ukrainian pa-tent «Plasmatron for spraying of coatings»No. 2002076032) UA, B23K10/00.

Design and operating parameters of the mi-croplasmatron provide formation of the laminarplasma jet (Reynolds criterion is 0.10-0.55). Ac-cording to this criterion, the microplasma spray-ing process is characterised [8] by:

• low thermal power, this making it possibleto decrease heating of the substrate and depositcoatings on small-size and thin-walled pieceswithout substantial local overheating and buck-ling;

• low level of noise in spraying with the lami-nar plasma jet, which is no more than 30—50 dB,this allowing avoidance of cumbersome protec-tion chambers;

• small size of the spraying spot (1—5 mm) ata small diameter of the nozzle equal to 1—2 mm.

The latter parameter is the key one for imple-mentation of the laser-microplasma claddingprocess, as it provides adequacy of the geometricsize of the spraying spot to the focal spot of thelaser. Therefore, with a spraying spot of about5 mm, it can be completely overlapped by thefocal spot of the laser, the thermal power densityin the spot being sufficient for remelting of thesprayed layer and its fusion with the substrate.

Flow diagram of the cladding process is shownin Figure 1. A specimen (plate) was mounted ona working frame approximately at equal anglesto axes of the laser beam and plasma jet. Thelaser beam was fed vertically from above. Theplasma jet transporting the cladding powder wasdirected to the focusing spot normal to the laserbeam. The laser beam and microplasma jet action

Table 1. Chemical composition of nickel alloy powders (Ni –base), wt.%

Powder grade Cr B Si Fe C

PG-12N-01 8—14 1.7—2.8 1.2—3.2 2—5 0.3—0.6

PG-12N-02 10—16 2.0—4.0 3.0—5.0 3—6 0.4—0.8

11/2012 17

zones were combined to form the common zone,the frame with the specimen mounted on it beingmoved relative to this common zone. Additionalscanning by using a scanator (Figure 2) was usedto smooth down the surface of the depositedlayer.

The principle of operation of this device is asfollows: DC motor 3 (see Figure 2, b), the rota-tion frequency of which is adjustable within 10—200 rpm, imparts the torque moment to eccentric2, whose axis is shifted relative to the motor axis.The eccentric engages holder 1, thus forcing itto oscillate relative to the semi-axes. The lensfixed in the holder oscillates together with it.This leads to periodic deflections of the radiationfocusing axis from the vertical position, whichin turn shifts position of the focal spot. As aresult, the laser beam focusing spot on a work-piece starts oscillating at a certain frequency de-pending on the rotation frequency of motor 3.The amplitude of such oscillations depends onthe value of eccentricity, which is set by using

eccentric 2. Return of the holder back to theinitial position is provided by spring 5, which isconstantly kept in the compressed state. The en-tire structure is mounted on swinging bracket 4,which makes it possible to arbitrarily select di-rection of the oscillations relative to the lasertreatment direction. This allows both transverseand longitudinal oscillations of the beam.

Experimental. To determine dependence ofheight h (mm) of the deposited layer on the proc-ess parameters, initially the process was per-formed by depositing single beads on a plate ofsteel St3 (δ = 8 mm). The following parameterswere chosen as the variable ones: laser radiationpower Plaser (kW), energy input E (J/mm), andspecimen movement speed v (m/h). Powder con-sumption Gp during the experiments was variedwithin 0.1—0.2 g/s. Other process parameterswere kept constant: diameter of the spot of thebeam focused on the specimen surface dsp = 5—6 mm, plasmatron current I = 43 A, voltage U == 30 V, plasma gas (argon) flow rate Q == 80 l/h, and shielding gas (argon) flow rateQsh = 240 l/h. To optimise the value of overlap-ping of the beads (according to the criterion ofroughness of the resulting coating on a similarplate), several beads were deposited by overlap-ping 10—50 % of their width.

Decrease in height of the deposited bead withincrease in power of the laser beam (Table 2,specimen 1) is related to a burn-off loss of partof the cladding material, as well as to overheatingof the base metal and dissolution of part of thedeposited bead material in it. Increase in heightof the bead (Table 2, specimen 5) is attributableto the noted instability in feeding the powder.

Results and discussions. Based on the experi-mental results, the mode corresponding to speci-

Figure 1. Flow diagram of the combined laser-microplasmacladding process: 1 – frame; 2 – specimen; 3 – laserbeam; 4 – plasma jet; 5 – microplasmatron

Figure 2. Appearance (a) and schematic of structure of the laser radiation scanator (for designations see the text)

18 11/2012

men 4, which provided the defect-free layers ata comparatively low energy input, was chosen toimplement the process of laser-microplasma clad-ding of steels with powders of the Ni—Cr—B—Sisystem alloys. Decrease in energy input led toformation of such a defect as microcracks (seeTable 2, specimens 1 and 5). Further investiga-tions showed that with increase in the powderconsumption to Gp = 0.5—0.8 g/s the speed ofmovement can be increased to 60 m/h, otherparameters being kept unchanged. This will pro-vide the deposited beads with the geometry simi-lar to that described in Table 2, along with de-crease in the HAZ. In addition, it was establishedthat the acceptable roughness of the depositedlayers (about Ra = 200—300 μm) occurs at thecoefficient of overlapping of the beads equal toKov = 25—30 %. It means that at a bead width of

6 mm the transverse movement of a specimen fordeposition of each next bead will be not less than4 mm. Also, it was established that the optimalparameters for laser-microplasma cladding canbe provided at an energy input ranging from 500to 800 J/mm. For comparison, it should be notedthat in laser powder cladding the energy inputis 120—250 J/mm [9]. This shows that overheat-ing of the deposited layers and increase in sizeof HAZ should take place in the case of combinedcladding, in contrast to laser cladding.

Overheating of the layers deposited by thelaser-microplasma method leads to some decreasein their hardness. The higher the energy input inthe process and, hence, the higher the tempera-ture in the working zone, the bigger this decreaseis. This is explained by the fact that in plasmaspraying of self-fluxing alloys at a temperature

Table 2. Effect of power and speed of movement of a specimen in CLMPC on height and quality of the bead deposited with powdersPG-12N-02

Specimennumber

Plaser, kW v, m/h E, J/mm Microstructure, ×20 h, mm Note

1 4 30 635 0.4 Presence of microcracks

2 3 6.5 2380 1.2 Bead is non-uniform inheight

3 3 10 1545 0.8 Bead is uniform

4 3 20 770 0.6 Same

5 3 30 520 1.0 »

6 2 10 1185 0.3 Bead is non-uniform inprofile

7 2 20 590 — Bead is not formed

11/2012 19

close to their melting temperature the burn-off(oxidation) of boron takes place to form B2O3.According to the data of study [10], at a tem-perature of 2000 °C the content of B2O3 in theoxide film within the process zone amounts to81 at.%. Under the combined laser-microplasmaprocess conditions the temperature of the work-ing zone is higher than in plasma spraying, thisintensifying burning-off of boron.

As shown by the experiments, in depositionof the layers not less than 0.6 mm high at anenergy input ranging from 300 to 400 J/mm thedecrease in hardness is minimal. The layers de-posited at the above energy inputs have hardnessthat corresponds to the certificate hardness ofthe applied cladding alloys. Investigations of thelayers produced by the CLMPC method showedthat on all the specimens the deposited layershave a sufficiently fine cast structure. Moreover,formation of columnar dendrites of metal, whichgrew in a direction of heat removal from the zoneof fusion with the base metal, takes place in thelower part of the deposited layers. In the upperpart of the deposited layers, the columnar den-drites propagate, as a rule, into the zone of finerequiaxed crystals, which is accompanied by someincrease in microhardness. The microhardness ofthe layers deposited at the specimen movementspeeds of up to 30 m/h in most cases amountsto about 3000 MPa.

Examinations of structures of the depositedspecimens showed the following. Specimen 1 (seeTable 2 and Figure 3, a) differed from the restof the specimens in the presence of structuraldefects, such as transverse cracks in the cast struc-

ture that propagated along the boundaries of den-drites in the deposited metal. HAZ in the basemetal was rather big. Its width was 2.5 timesbigger than cladding thickness. Specimens 2—4(see Figure 3, b—d) had no cracks and no sepa-rations from the substrate. Microcracks similarto those observed in specimen 1 were detected inspecimen 5, which can be explained by close en-ergy inputs in cladding of these specimens. Speci-mens 6 and 7 sprayed at a lower power of thelaser were characterised by formation of a lower-quality bead. In all the specimens the depositedmetal had the cast dendritic structure, transform-ing into the fine-crystalline one in the upper part.Interlayers consisting of nickel borides Ni3B andnickel silicides Ni2Si and, probably, their eutec-tics with γ-Ni, as well as chromium carbidesCr23C6 and Cr7C3, were located along theboundaries of light dendrites, which were γ-Nibased solid solution.

The feature in common to the specimens isthat the dendritic structure of the cladding nearthe zone of fusion with the base metal was freefrom inclusion. Structure of the fusion region(white strip) consisted of γ-Ni solid solution andhad a decreased hardness, 25 % lower, on theaverage, than hardness of the cladding. The re-gion located below the fusion zone (HAZ) canbe subdivided into two parts as to its depth: aregion adjoining the fusion zone, having hardnessof 2590—3260 MPa, and a region located below,which adjoins the base metal and has hardnessof 1580—1940 MPa. Microhardness of the basemetal was 2100—2310 MPa, on the average. Pre-sumably, the presence of the HAZ metal regions

Figure 3. Microstructures (×100) of the layers deposited by the CLMPC method using powder PG-12N-02: a—f –specimens 1—6, respectively, from Table 2

20 11/2012

with a different hardness can be explained by adiffusion redistribution of elements. Probably,such alloying elements as carbon, boron and sili-con propagated from the deposited layer into thatpart of HAZ which adjoins the transition part.Moreover, it is likely that carbon from the lowerpart of HAZ redistributed to the upper part. Sizeand hardness of the cladding regions and metaldepend on the combination of parameters of thelaser and microplasma processes, consumption ofthe additive powder and speed of movement ofa specimen in CLMPC.

The experiments showed that increase in thelaser beam focusing spot up to values of dsp == 5—6 mm leads to the need to use a substantialpower of laser radiation (about 3 kW). To de-crease the latter and reduce roughness of the sur-face of the layer, the laser beam was additionallyscanned across the cladding with amplitude of2 mm and frequency of about 20 Hz. The beamwas scanned by using a scanator (see Figure 2).

Adding of scanning of the laser beam acrossthe CLMPC direction allowed diameter of thespot focused on the surface treated to be de-creased to 4 mm, and laser radiation power to bereduced to 2 kW. Other process parameters werekept constant. Adding of scanning of the laserbeam reduced the sensitivity of the deposited lay-ers to cracking. The general trends in formationof structure in these layers remained unchanged(Figure 4, a, b).

Quality of the resulting layers also dependson the consumption of the additive powder. Forexample, the CLMPC process without scanningof the laser beam, at the PG-12N-02 powder con-sumption of Gp = 1.0—1.2 g/s, allowed deposit-ing the sound layers 0.5—0.6 mm high at a speedof 50 m/h and at radiation power P = 3 kW. Inthis case the size of HAZ was approximately equalto height of the deposited coating (Figure 4, c).Therefore, the indicated consumption of the ad-ditive powder materials for CLMPC is 0.8—1.2 g/s.

The results obtained in the above experimentswere compared with the results of cladding of

similar materials performed by the laser powdercladding method developed at the E.O. PatonElectric Welding Institute [9]. It turned out thatthey were rather close in value of irregularities(roughness) and appearance of the clad surfaces.The main difference consisted in sticking of aninsignificant quantity of the powder material tothe surface in CLMPC.

It was found that the layers of the Ni—Cr—B—Sisystem alloys deposited by the laser method hada cracking index of about 40—60 % (α = 0.4—0.6),whereas the combined cladding allowed decreas-ing this index from 10—20 % (α = 0.1—0.2) to acomplete elimination of microcracks.

Both standard friction machine 2070 SMT-1and in-house friction machine were used to de-termine wear resistance of the deposited layers.

Figure 4. Microstructures of the layers deposited by the CLMPC method using powder PG-12N-02 at a process speed of20 (a, b) and 50 (c) m/h with transverse scanning of the laser beam (a – ×25; b – ×100) and without it (c – ×32)

Figure 5. Comparison of wear resistance and hardness HRCof Ni—Cr—B—Si system alloys deposited by different methodswith those of steel 38KhN3MFA in dry friction: 1 –38KhN3MFA; 2 – plasma spraying with powder PG-12N-02; 3, 4 – laser cladding with PG-12N-01 and PG-12N-02,respectively; 5, 6 – combined deposition of layers withPG-12N-01 and PG-12N-02

11/2012 21

Wear resistance in dry sliding friction was de-termined in percent, by taking wear resistanceof steel 38KhN3MFA with hardness HRC 42—43as 100 %. It was established as result that thewear resistance values in CLMPC can be superiorto those characteristic of laser cladding. How-ever, in a case of failure to comply with thethermal conditions, i.e. overheating of specimensat low cladding speeds, the value of wear resis-tance may decrease to a substantial degree be-cause of weakening of the hard phases. An ex-ample of such decrease in wear resistance of metaldeposited with powder PG-12N-01 is shown inFigure 5. There this indicator decreased almostto 60 % relative to the same indicator for steel38KhN3MFA.

CONCLUSIONS

1. The efficiency of applying CLMPC is deter-mined by decrease in the quantity of microcracksin the deposited layers. For instance, the layersof Ni—Cr—B—Si system alloys (PG-12N-01 andPG-12N-02) deposited by the laser method hada cracking index of about 40—60 %, whereas com-bined cladding of the same alloys allowed de-creasing this index from 10—20 % to completeelimination of microcracks.

2. Comparative dry friction tests of specimensof the base metal (steel 38KhN3MFA, whosewear resistance was taken as 100 %) and speci-mens deposited with the same Ni—Cr—B—Si sys-tem alloys showed the possibility of providingwear resistance of an order of 120—130 % in laserpowder cladding, and more than 140 % inCLMPC.

3. Along with the above advantages, CLMPChas certain drawbacks compared to laser clad-ding. The key drawbacks include increase in sizeof HAZ in the base metal, decrease in hardnessof the deposited layers as a result of weakeningof metal (burn-off loss (oxidation) of boron and,hence, decrease in the content of the boridephases, as well as coagulation of particles of the

strengthening carbide and silicide phases). Thecause is increase in temperature of the workingzone due to a substantial growth of the processenergy input (to 500—800 J/mm), compared tolaser cladding (normally 120—250 J/mm), whichis required to achieve the optimal parameters.

4. Further investigations on elimination of thesaid drawbacks will show expediency of applyinglaser-microplasma cladding for deposition ofwear-resistant coatings both in manufacture andin repair of parts of the shaft type operating infriction pairs (e.g. components of sleeve assemblyof internal combustion engines, and running gearof motor and railway transport).

1. Borisov, Yu.S., Kharlamov, Yu.A., Sidorenko, S.L.et al. (1987) Thermal spray coatings from powdermaterials: Refer. Book. Kiev: Naukova Dumka.

2. Grigoriants, A.G., Safonov, A.N., Shibaev, V.V.(1982) Production of wear-resistant chrome-nickeland chrome-boron-nickel coatings by using laser ra-diation. Izvestiya Vuzov. Mashinostroenie, 3, 87—92.

3. Safonov, A.N., Grigoriants, A.G., Shibaev, V.V. etal. (1984) Investigation of crack formation in lasercladding of chrome-boron-nickel powdered alloys.Ibid., 12, 64—68.

4. Dilthey, U., Wieschemann, A. (2000) Prospects bycombining and coupling laser beam and arc weldingprocesses. Rivista Italiana della Saldatura, 52(6),749—759.

5. Coddet, C., Montaron, G., Marchione, T. et al.(1998) Surface preparation and thermal spray in asingle step: the PROTAL process. In: Proc. of 15thITSC (Nice, France), Vol. 1, 1321—1325.

6. Krivtsun, I.V. (2002) Combined laser-arc processesof materials treatment and devices for their realiza-tion: Syn. of Thesis for Dr. of Techn. Sci. Degree.Kiev: PWI.

7. Shelyagin, V.D., Krivtsun, I.V., Borisov, Yu.S. etal. (2006) Laser-arc and laser-plasma welding andcoating technologies. Svarka v Sibiri, 1, 32—36.

8. Borisov, Yu.S., Vojnarovich, S.G., Bobrik, V.G. etal. (2000) Microplasma spraying of bioceramic coat-ings. The Paton Welding J., 12, 63—67.

9. Velichko, O.A., Avramchenko, P.F., Molchan, I.V.et al. (1990) Laser powder cladding of cylindricalparts. Avtomatich. Svarka, 1, 59—65.

10. Gershenzon, S.M., Borisov, Yu.S., Razikov, M.I. etal. (1978) Interaction of self-fluxing nickel alloyswith oxygen. Svarochn. Proizvodstvo, 1, 9—11.

22 11/2012

THERMODYNAMIC ANALYSIS OF SLAG MELTSIN MANUFACTURE OF FUSED WELDING FLUXES

L.A. ZHDANOV1, A.N. DUCHENKO1, I.A. GONCHAROV2, V.I. GALINICH2,A.V. ZALEVSKY2 and N.Ya. OSIPOV3

1NTUU «Kiev Polytechnic Institute», Kiev, Ukraine2E.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

3OJSC «Zaporozhstekloflyus», Zaporozhie, Ukraine

Applicability of equilibrium thermodynamic laws for analysis of physico-chemical reactions was establishedproceeding from analysis of temperature-time conditions of melting of welding fluxes in arc and open gasfurnaces. Thermodynamic analysis of reactions of charge components used in welding flux manufacture wasconducted. Techniques to control the process of slag melt refining are determined.

Keywo r d s : fused welding fluxes, slag melt, sul-phur, phosphorus, thermodynamic analysis

Ukraine today is a leading manufacturer of fusedwelding fluxes. Such leading scientists as E.O.Paton, K.V. Lyubavsky, B.E. Paton, V.I. Dyat-lov, I.I. Frumin, V.V. Podgaetsky, I.K. Pokhod-nya and many others participated in developmentof flux compositions and their manufacturingtechnology. In connection with depletion of de-posits of mineral raw materials traditionally ap-plied in fused flux manufacture, the problem ofensuring the required flux composition and theirquality, respectively, became much more acute.

Conducted analysis of charge materialsshowed that in most of the cases the content ofimpurities in them is specified by the normativedocuments (GOST, DSTU, TU). Technicaldocumentation predominantly specifies the con-tent of the main component in the raw material,and in a number of cases does not limit the contentof impurities – sulphur, phosphorus, and ironoxides. At the same time, their content in thefluxes is limited. On the other hand, statisticaldata of incoming inspection of the raw materialsare indicative of wide ranges of variation of theirimpurity content.

The most contaminated are manganese ore andfluorospar concentrates, in which sulphur andphosphorus content reaches 0.3 % in some cases.These materials make up almost half of the chargein manufacture of the most widely acceptedfluxes of AN-348-A, OSTs-45 grades. Therefore,it is possible to keep their sulphur and phosphoruscontent only at the upper admissible level by thetechnical requirements. Forced application oflow-grade ore materials leads to increase of theamount of impurities contributed by them to the

melt that in its turn reduces the technologicalmargin on impurities in flux melting.

Therefore, work on investigation of processesrunning in flux-meting furnaces is urgent for de-velopment of recommendations on lowering theimpurities in the slag melt.

Fused welding fluxes are made in open gasand arc furnaces [1]. These melting units differby temperature conditions, lining type, volume,mixing conditions and time of slag melt exist-ence. For open gas furnaces these are: up to1450 °C temperature, up to 60 t melt volume,and up to 6 h melting time. In the case of arcfurnaces slag melt temperature is higher on av-erage and can reach 1800—1900 °C. Melt volumefor various types of furnaces can be in the rangefrom 50 kg up to 5 t, and melting time is from 1up to 2 h. Intensive processes of slag melt mixingproceed in arc furnaces under the impact of themagnetic field and temperature gradient. Thus,equilibrium conditions are in place in both thecases, which are characterized by long-time exist-ence of the melt, large volume and uniformity inlocal melting regions. All that allows applying theprinciples and laws of chemical equilibrium ther-modynamics for assessment of physico-chemicalprocesses in flux-melting furnaces.

In the general case the melting space of afurnace can be regarded as a closed thermody-namic system that is related to features of massexchange with the environment. In flux meltingby a traditional schematic, uniformly blendedcharge is fed into the furnace and, as a rule, thereis no further adding of charge components to thefurnace volume. Mass transfer can be performedonly as a result of gas removal from the meltingspace and transition of compounds from slag intothe metal phase.

© L.A. ZHDANOV, A.N. DUCHENKO, I.A. GONCHAROV, V.I. GALINICH, A.V. ZALEVSKY and N.Ya. OSIPOV, 2012

11/2012 23

Melting space, in which charge material com-ponents are present in the form of solid and par-tially molten particles of the slag melt, shot ironand lining, is a heterogeneous system. To performanalysis of such a complex heterogeneous ther-modynamic system, the melting space should beseparated into certain homogeneous systems(zones), in which chemical reactions will be con-sidered. Zone interaction is performed throughinterfaces. It is understandable that such a divi-sion is conditional, and does not reflect all thediversity of high-temperature processes in the fur-nace melting space, in particular, kinetics of slagmelt homogenizing, hydrodynamic conditions ofits existence, process of gas removal from themelt and impact of the electromagnetic field (influx melting in electric arc furnaces).

Traditionally [2] the process of flux meltingis divided into three stages: reaction in the solidstate, flux formation and slag melt homogeniz-ing. As a result, the following phases can besingled out in the flux volume:

• solid, in which thermal dissociation and in-teraction of charge particles take place;

• partially molten dispersed solid-liquidphase, arising as a result of charge heating, eu-tectic and contact melting;

• slag melt with a certain degree of homoge-nizing;

• metallic phase (shot iron) forming as a resultof chemical reactions in the slag melt;

• gas phase forming as a result of thermaldissociation of charge components and chemicalreactions with formation of gaseous products,which float to the slag melt surface in the formof bubbles.

Interfaces are well developed, rather blurred,particularly in the presence of dispersed particlesat the beginning of slag melt formation. A certaintemperature interval corresponds to each phasein flux manufacture.

The objective of this work was analysis ofthermodynamic probability of chemical reactionrunning in the flux-melting furnace volume, andprecising the mechanism of sulphur and phospho-rus removal from the slag melt to lower the con-tent of these impurities in the finished flux com-position. A characteristic of the probability ofreaction running was dependence of Gibbs energyΔG on temperature. In metallurgy in most of thecases a simplified Gibbs equation is used, whichallows for the change of enthalpy and entropyof the reaction, depending on temperature. Heatcapacity of initial materials and reaction productsin this case is neglected. This is related to thefact that heat capacity contribution to ΔG value

at up to 800—1000 K temperatures is negligible.At temperature rise heat capacity value rises bya logarithmic dependence, and it can change theheat capacity of elements 2 times. One of thecalculation methods of allowing for the changeof heat capacity value, depending on tempera-ture, is application of Uhlich function

M0 = ln T

298.15 +

298.15T

— 1. (1)

As a result, equation for calculation of thechange of Gibbs energy becomes

ΔGT0

= ΔH2980

— TΔS2980

— ΔCp2980

TM0, (2)

where

ΔH2980 = ∑ H298 prod

0 — ∑ H298 in0 ; (3)

ΔS2980 = ∑ S298 prod

0 — ∑ S298 in0 ; (4)

ΔCp2980 = ∑ ΔCp298 prod

0 — ∑ ΔCp298 in0 ; (5)

ΔH2980 , ΔS298

0 , ΔCp2980 is the variation of values

of thermodynamic characteristic functions of theparticipants (products and initial materials) ofthe chemical reaction under standard thermody-namic conditions (at temperature of 298 K andatmospheric pressure of 1 atm (9.80665⋅104 Pa)).

One of the main problems, arising during ther-modynamic metallurgical calculations, is findingthe values of enthalpy, entropy and heat capacityof the initial materials and reaction products.Unfortunately, the most fundamental works [3,4] do not include the data on complex com-pounds, so that we assumed them on the basis ofprocessing the experimental equations, used inmetallurgical calculations [5].

In manufacture of fused welding fluxes forsteel welding, the most often used raw materialsare manganese ore concentrates, quartz sand, alu-mina, periclase powders, lime, marble, fluorite,fluorspar, rutile, zirconium concentrates, etc.,which contain such chemical compounds as SiO2,MnO2, Mn2O3, MnO, Fe2O3, CaO3, MgCO3,TiO2, FeS2, MnS, MnnP, CaF2, P2O5,Ca5(PO4)3(F, Cl, OH). In ores phosphorus ismostly present in the form of phosphorus-calciumsalt, included into the composition of apatiteminerals [6]. In addition, material of the elec-trodes and lining – carbon (for arc furnaces)and firebrick (for open gas furnaces) – willparticipate in the interaction reactions.

At consideration of the first stage, analysis ofchemical reactions in the solid state – gaseouscompounds formation and removal – is tradi-tionally performed. This is exactly the stage atwhich melting of the flux charge proceeds, i.e.charge transition from the solid into the liquid

24 11/2012

state. The charge, which later on forms the slagmelt, is first present in the form of separate com-ponents, the melting temperature of which ismuch higher than that of the melting space. Thecharge melting process proceeds at the expenseof contact melting of charge particles with si-multaneous thermal dissociation of carbonatesand higher oxides.

In the contacting surfaces of flux charge par-ticles interdiffusion takes place, which results information of a eutectic layer and contact meltingat temperatures below the melting temperatureof individual charge materials. This results information of partially molten dispersed phase,which promotes further charge melting at melt-ing space temperatures, and flux forming processtakes place. The lowest-melting eutectics canform as a result of contact interaction of fluxcharge particles already at temperatures of 900—1100 °C [7, 8].

Processes of thermal dissociation of chargecomponents are accompanied by intensive gasevolution, thus leading to increase of charge melt-ing rate due to mixing of the forming liquidphase. It should be noted that formation of gase-ous compounds should influence the kinetic fea-tures of charge material interaction, and canchange the concentrational conditions of the re-actions between the components in the solid andsemi-liquid state.

At the first stage reactions of carbonate de-composition with carbon dioxide gas evolution,as well as reduction of higher manganese oxides(Figure 1, a) with formation of gaseous oxygen,take place. Manganese, calcium and magnesiumsulphides do not decompose (Figure 1, b). Inter-action of oxygen with sulphides results in ap-pearance of gaseous sulphur oxide SO2, which isremoved from the melting space (Figure 2, a).Therefore, it is believed that for maximum re-moval of sulphur at the first stage of melting itis necessary to create oxidizing conditions, whichcan be formed due to dissociation of higher oxides,for instance MnO2 (see Figure 1, a). Besides oxy-gen, sulphur can be removed from the compoundsas a result of interaction of calcium sulphide withhigher manganese oxide (Figure 2, a).

Manganese phosphides can decompose withformation of solid and gaseous phosphorus (seeFigure 1, b). The thus formed manganese canhere interact with phosphorus oxide also withformation of solid and gaseous phosphorus, whichcan also be reduced by other metals, for instance,silicon (Figure 2, b). However, their presence atthe first stage of melting is improbable. Reactionof phosphorus oxidation at interaction with sili-con oxide is more probable. Reduction of phos-

phorus oxide by iron at this melting stage is im-probable (see Figure 2, b).

Proceeding from the calculation data, it canbe anticipated that phosphorus, similar to sul-phur, should be removed already at the first stageof melting. This process, however, is preventedby two interrelated factors: under oxidizing con-ditions sulphur is removed at the first meltingstage during a reaction of interaction of solid andgaseous phosphorus with formation of P2O5,which, in its turn, comes into a reaction withcalcium and magnesium oxides, forming the re-spective phosphates (Figure 3, a). At the same

Figure 1. Change of Gibbs energy for reactions of dissocia-tion of carbonates, higher oxides of manganese (a),phosphides of manganese and sulphides (b)

Figure 2. Change of Gibbs energy for reactions of formationof sulphur oxide (a) and phosphorus precipitation (b)

11/2012 25

time, these oxides can form complex compoundsbased on silicon oxide. The known property ofsilicon oxide to form silicon-oxygen chains in themelt leads to phosphorus oxide being bound intoextremely strong aggregated complex com-pounds, which further on at high contents ofsilicon oxide in the melt are built into the sili-con-oxygen lattice of the slag melt and can beremoved from the melt only when reducing condi-tions are created. The probability of these reactionsis higher (Figure 3, b). Temperature, at which thismelting stage is over, can be conditionally takento be equal to about 1000—1200 °C (±50 °C). Theprocess of charge melting is determined by thekinetics of the processes of gas evolution and con-tact melting of charge particles. It is exactly thekinetics of these processes at this stage that deter-mines the completeness of removal of gaseous prod-ucts from the melt, namely phosphorus and sulphuroxides. The main factor, determining the comple-tion of this melting stage, is removal of oxygenand carbon dioxide gas from the melt.

The flux forming (slag forming) stage is char-acterized by that all the charge mass turns intothe melt as a result of interaction of eutecticsformed at contact melting with the charge bulk.Charge homogeneity is not achieved [2]. Its masscontains a large number of gas bubbles and hasa non-uniform structure. At traditional conduct-ing of the melting process, formation of thermo-dynamically stable complex compounds – cal-cium, magnesium, manganese, iron and phospho-rus silicates goes on at this stage in the presenceof silicon oxide in the charge (see Figure 3, b).Calculation results show that due to known prop-erty of silicon oxide, mentioned above reactions

proceed in the entire temperature range charac-teristic for the flux melt, i.e. complex formationaccompanies the entire process of flux melting.These are exactly the complex compounds thatdo not allow phosphorus to be removed from theflux melt in the process of flux melting. Note thefact that the probability of CaO⋅P2O5 formationrises with temperature, thus making phosphorusremoval from the slag melt difficult.

The objective of phosphorus removal from themelt can be implemented by creation of certainconcentrational conditions, in which complexformation will be limited, for instance, by addi-tion of silicon oxide (or part of it) separatelyfrom the other charge components.

The main outcome of the second stage, whichcorresponds to temperature interval of 1200—1270 °C (±50 °C), is charge transformation intothe slag melt, in which unmolten charge particlesare absent.

At the third stage the processes of slag melthomogenizing and degassing proceed. An impor-tant requirement to oxide fluxes for steel weldingis lowering of their oxidizing ability relative tometal in the reaction welding zone. Therefore,during their manufacture the slag melt is furtheroxidized through reactions of carbo-thermal re-duction of the melt components with the carbonof the lining (in the case of arc furnaces) or cokebreeze additives (in the case of open gas fur-naces). As a result, oxidizing conditions in thefurnace are replaced by reducing conditions, fluxcomponents interact with carbon, forming themetal phase, containing iron, manganese and sili-con (Figure 4). Appearing metals interact withphosphorus oxide with formation of gaseous andsolid phosphorus (see Figure 2, b), taking it outof the slag melt.

Thus, the most important in terms of slag meltdephosphorization is the reaction of phosphorustransfer into the metal phase, which is realizedexactly at this stage of the process. Phosphorusreduced by metals goes into the metal phase to-gether with other metals and precipitates on thefurnace bottom plate. Temperature and concen-trational conditions of this process running havethe main role here. On the one hand, reactionsof intermediate manganese oxides of Mn2O3 typeare highly probable, and on the other – prob-ability of the reactions of iron oxide reductionexceeds the possibility of MnO reduction. Thisenables controlling the processes of manganeseloss as a result of its transition into the metalphase. In particular, such a lowering can beachieved due to complete transfer of higher ox-ides into the lower (MnO) oxide at the first andsecond melting stages.

In addition, carbon reduces phosphorus fromcomplex compounds based on calcium and man-

Figure 3. Change of Gibbs energy for reactions of formationof phosphorus oxide and complex compounds with phos-phorus oxide (a) and silicon oxide (b)

26 11/2012

ganese oxides in the presence of silicon oxide(Figure 5). However, some of these reactionsstart running only after 1200 °C, and their bulk –after 1500 °C, i.e. at the final stage of melting.The thus formed phosphorus precipitates on fur-nace bottom. In arc furnaces the probability of thereactions is enhanced in near-electrode regions. Thesimultaneously precipitating carbon oxide pro-motes homogenizing of the slag melt.

Completeness of phosphorus oxide removalfrom the slag melt is determined by the presenceof the metal phase. Melting conditions can beselected so that it will mainly consist of iron,whereas silicon and manganese transition will beminimum. Temperature range of the last meltingstage is determined by technological capabilitiesof arc furnaces. For open gas furnaces it is overat the temperature of 1470 °C (±10 °C), and forflux melting – at 1800—1900 °C.

CONCLUSIONS

1. Proceeding from analysis of temperature-timeconditions of welding flux melting in arc andopen gas furnaces, the possibility of applicationof the laws of equilibrium thermodynamics foranalysis of physico-chemical reactions was deter-mined. Here the heat capacity of the elementsand their compounds should be taken into ac-count, and the calculation proper should be per-formed by Uhlich formula.

2. As a result of thermodynamic analysis ofthe reactions of charge components used in weld-ing flux manufacture it is established that:

• reactions of sulphide decomposition do notproceed, and their removal requires the presenceof oxidizing conditions;

• reactions of phosphide decomposition run inthe entire temperature range, but under oxidizingconditions phosphorus oxide forms, which goesinto difficult to remove complex compounds;

• in the presence of carbon, exchange reactionsproceed in the melt, which result in formationof metals reducing phosphorus from the oxide upto the metallic and gaseous state, and oxides ofthe same metals, i.e. this group of reactions are

interrelated and should be regarded as one ther-modynamic system;

• as a result of exchange reactions with par-ticipation of carbon, silicon oxide and phos-phates, which are present in complex compounds,reactions of phosphorus reduction with simulta-neous formation of carbon oxide run at the thirdstage of melting. There are ten CO molecules forone P4 molecule that should promote removal ofgaseous phosphorus from the melt. Carbon mono-xide further oxidizes in the slag melt, promotinglowering of oxidizing ability of the ready flux.

3. Methods to control the processes of slagmelt refining are as follows:

• creating oxidizing conditions at the initialmelting stage by adding higher oxides of variablevalency for sulphur transfer into gaseous oxides;

• separate addition of charge components, inparticular silicon oxide, separately from the restof the charge bulk to prevent formation of com-plex compounds in the slag melt, hindering phos-phorus removal;

• mandatory simultaneous addition of carbonand silicon oxide at the final stage of meltingprocess for decomposition of phosphates presentin the complex compounds;

• mandatory presence of the metal phase,forming through reactions of reduction of iron,manganese and silicon oxides by carbon, to re-move solid phosphorus from the slag melt. It ispossible to create such concentration and tem-perature conditions, under which the metal phasewill consist mainly of iron and phosphorus, andmanganese transition into it will be minimum.

1. Podgaetsky, V.V., Lyuborets, I.I. (1984) Weldingfluxes. Kiev: Tekhnika.

2. Podgaetsky, V.V. (1947) Production of AN-3 fluxfor automatic welding. Kiev.

3. (1965—1982) Thermal constants of materials: Refer.Book. Ed. by V.P. Glushko. Moscow: VINITI.

4. (1978) Thermodynamic properties of individual ma-terials. Ed. by V.P. Glushko. Moscow: Nauka.

5. Kazachkov, E.A. (1988) Calculations on theory ofmetallurgical processes. Moscow: Metallurgiya.

6. Leontiev, L.I., Yusfin, Yu.S., Malysheva, T.Ya. etal. (2007) Raw and fuel materials base of iron in-dustry. Moscow: Akademkniga.

7. Zalkin, V.M. (1987) Nature of eutectic alloys andeffect of contact melting. Moscow: Metallurgiya.

8. Savin, V.S., Mikhalyova, O.V., Zubova, Yu.A. (2007)Diffusion of atoms from liquid to solid phase in contactmelting. Zhurnal Tekhn. Fiziki, 33(10), 27—32.

Figure 4. Change of Gibbs energy for reactions of oxidereduction by carbon

Figure 5. Change of Gibbs energy for reduction of phos-phates by carbon and silicon oxide

11/2012 27

REPAIR WELDING OF INTERMEDIATE CASESOF AIRCRAFT ENGINES FROM HIGH-TEMPERATURE

MAGNESIUM ALLOY ML10 WITH APPLICATIONOF ELECTRODYNAMIC TREATMENT

L.M. LOBANOV1, N.A. PASHCHIN1, A.V. CHERKASHIN1, G.I. TKACHUK1, V.V. SAVITSKY1, O.L. MIKHODUJ1, K.V. SHIYAN1, V.K. LEVCHUK1, V.V. ZHYGINAS2 and A.P. LYASHCHENKO2

1E.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine2SC «Plant 410 of Civil Aviation», Kiev, Ukraine

Technology was developed for repair welding of damages in aircraft engine intermediate cases from mag-nesium alloy ML10. The technology comprises electrodynamic treatment of welds aimed at reducing thelevel of residual welding stresses. It was experimentally proved that treatment practically eliminates residualstresses in the weld. At charging voltage of up to 200 V electrodynamic treatment operator can performmaximum 1100 electrodynamic impact operations per shift, and at 500 V voltage – not more than 100operations, that fully meets the requirements of production cycle of repair welding of aircraft intermediatecase.

Keywo r d s : argon-arc repair welding, electrody-namic treatment, magnesium alloy, aircraft engine cases,magnetic field intensity, pulsed current, charging volt-age, capacitor capacitance, welding stresses, treatmenteffectiveness

Development of modern technologies of repairof aeronautical engineering equipment is relatedto searching for new ways of extension of servicelife of metal structures from high-temperaturemagnesium alloys, reconditioned by repair weld-ing. One of the causes for shortening of the serv-ice life of flying vehicles are residual weldingstresses in repair welds, which adversely affectthe fatigue strength, corrosion resistance and re-sidual distortion of aircraft structural elements.This necessitates investigation of advanced meth-ods to control the stressed state of welded joints,one of which is treatment by electric currentpulses [1, 2].

Method of realization of pulsed current impacton metals is electrodynamic treatment (EDT)based on initiation of electrodynamic forces inthe material, arising at passage of a current dis-charge in the treated material [3]. The mecha-nisms of EDT impact on the treated material aredescribed in detail in [4].

One of the structural components of the air-craft, in which the damage is repaired by welding,is the aircraft engine intermediate case (AEIC).AEIC purpose is aircraft engine fastening on theaircraft wing and thermal insulation of the air-frame structural components from thermal im-

pact of an operating engine. Figure 1, a, showsAEIC appearance as-assembled with D-36 en-gine. Conditions of AEIC operation make highrequirements to fatigue and static strength char-acteristics of the structure at high (up to 400 °C)temperatures, as well as to its dimensional sta-bility, determining the aerodynamic and propul-sion performance characteristics of D-36 engine.Proceeding from that, static and fatigue strengthof AEIC repair welded joints should correspondto mechanical characteristics of base metal, andlevel of residual welding stresses – to minimumvalues. Thus, it is believed to be reasonable toassess EDT capabilities to lower the level of re-sidual welding stresses in AEIC repair welds.

The objective of this work is development ofthe technology of repair welding of AEIC damagewith EDT application.

AEIC is a large-sized cast structure from mag-nesium alloy ML10 (Figure 1, b) which consistsof outer 1 and inner 2 cylindrical shells, con-nected by stiffeners – posts 4. One of the designfeatures of the posts is presence of inner cavitiesin them, through which the coolant circulates,which is designed for minimizing the thermalimpact of operating engine on AEIC. Outer shellis designed for mounting AEIC on aircraft wing,and the inner shell – for fastening the aircraftengine 3.

The most characteristic damages of AEIC(Figure 2) rectified by repair welding, are fatiguecracks, disturbing the integrity of the post in the

© L.M. LOBANOV, N.A. PASHCHIN, A.V. CHERKASHIN, G.I. TKACHUK, V.V. SAVITSKY, O.L. MIKHODUJ, K.V. SHIYAN, V.K. LEVCHUK, V.V. ZHYGINAS and A.P. LYASHCHENKO, 2012

28 11/2012

points of their connection to outer and innershells (Figure 2, weld 1, view A). Formation offatigue cracks on the external surface of outershell in the zone of reinforcement for coolingpipeline flange (Figure 2, weld 2, sectional viewA—A) and on reinforcement for the system ofAEIC fastening to the wing is less common. Interms of design the assembly of AEIC fasteningto the wing is similar to that shown in sectionalview A—A. The result of mentioned service defectsare a partial loss of the load-carrying capacity ofthe structure and violation of leak-tightness ofAEIC cooling cavities.

Repair of AEIC damages was performed usingmanual single- and multipass nonconsumableelectrode arc welding (TIG) in shielding gas at-mosphere (argon) in the following modes: Ua == 20 V, vw = 1.5 mm/s. Shielding gas was pureargon of grade A, recommended for welding tightjoints, to which welds 1 and 2 belong (argonflow rate was 0.25—0.30 l/s). Post repair (seeFigure 2, weld 1) was performed at current of200—350 A in five passes, repair of reinforcementfor the cooling main pipeline (Figure 2, weld2) – at current of 200—250 A in two passes.Joint preparation for welding was performed bymechanical cleaning of the repair joint to thewidth of 15—30 mm from both sides using a steelbrush (stainless steel diameter of 0.2 mm) andscraping. Time interval between mechanicalcleaning and welding did not exceed 24 h. Fillerrods of ML9 grade of 6 mm diameter were used,the surface of which was treated by chemicaletching before welding. Preparation of crackedges was performed with the angle of openingof 50—70°, with more than 3 mm radius of openingin the root up to residual thickness of 0.3—0.5 mm. TIG welding was performed with con-current local preheating of the welding zone,which was realized by placing specialized heatersbased on tubular electric heating elements on thebase metal. Heating temperature was equal to150—200 °C. The first pass was made at minimumcurrent with the initial and final sections of therepair weld reaching the base metal. Here smoothtransition of the deposited to base metal wasensured with welding up of the crater in the modeof smooth extinction of the arc. At forced stop-ping of the welding process, because of filler rodreplacement, overlapping of earlier depositedweld by 20—30 mm was performed. The over-lapped surface was first cleaned mechanically.

Presence of residual stresses in AEIC repairwelds in a number of cases requires performanceof postweld heat treatment of the item in large-sized electric furnaces that is a highly energy-consuming operation. Application of heat treat-ment is required when repair welding of more

than two AEIC damages is performed. At thesame time, there are cases, when a unit defect ofsmall depth and length is to be repaired. Thenapplication of total heat treatment is not rational.Practical experience of application of postweldlocal heating of the repair weld with tubularelectric heating elements used for welding, dem-onstrated its low effectiveness as a result of highheat conductivity of ML10 alloy. Application of

Figure 1. Appearance of AEIC (1) as-assembled with D-36aircraft engine (2) (a) and AEIC (b—f for 1—4 see the text)

Figure 2. Schematic of location of repair welds at servicedamage of AEIC in the zone of connection of the post tothe outer and inner shells (weld 1) and in the zone offastening the cooling pipeline (weld 2)

11/2012 29

EDT will allow not only lowering the level ofresidual stresses in short repair welds withoutheat treatment application, but also replacing itin the future that will lower the cost of AEICreconditioning. It should be noted that by theresults of testing by static tension, EDT does nothave any negative influence on mechanical char-acteristics of AEIC repair welded joints.

EDT influence on distribution of residualstresses arising at two-pass deposition of weld 1was studied on samples of 350 × 200 × 8 mm size.Before bead deposition a cut of the length, widthand depth of 200, 1.6—2.0 and 8—10 mm, respec-tively, was made with a hand cutter along theweld by a procedure described above. In orderto reproduce the operations of AEIC repair, two-pass welding was performed in the cut section inthe mode given above. Here, the geometricalcharacteristics of the deposited weld corre-sponded to parameters of the repair joint madein AEIC on the shop floor.

EDT influence on the magnitude and distri-bution of residual stresses when making weld 2was studied on samples of 300 × 200 × 8 mm size,containing elements of reinforcement for theflange of the cooling pipeline, shown in Figure 3(sectional view A—A). Before deposition a cut ofthe length, width and depth of 50, 1.6—2.0 and8—10 mm, respectively, simulating the fracture,was made between the bosses, and its edge prepa-ration was made similar to weld 1. In order tosimulate repair welding performed at damage re-conditioning, single-pass deposit 50 mm long wasmade between the bosses, in the mode mentionedabove. After bead deposition and complete cool-ing of the samples, EDT of welded joints of thesamples was performed in the modes given inTable 1.

Welded joints were treated along the weldaxis in the direction from the middle towards theedges.

Before performance of TIG welding, evalu-ation of the initial level of stresses in ML10 alloywas performed by the method of electron speckle-interferometry on sample surface. After welding,values of longitudinal component σxx of residualstresses were determined in repair weld zone be-fore and after EDT performance. Treatment ef-fectiveness was assessed by the results of com-parison of stressed state parameters before andafter EDT.

Evaluation of initial stressed state on the sur-face of ML10 alloy samples before weldingshowed that stress distribution on their outersurface was uniform, while σxx values were inthe range of 4—6 MPa.

EDT of samples with deposited welds 1 and2 was performed by series of five current dis-charges in modes corresponding to charging volt-age U = 200 and 500 V. Sections on the surfaceof deposited beads were treated by applicationof current pulses with monitoring σxx variationin EDT zone. Initial and final weld sections of10 mm length, in which values of initial stressesare minimal, were treated in mode 1, and theother bead surfaces – in mode 2 from Table 1.

Initial σxx values in the metal of single-passweld 2 before and after treatment were equal to120 and 20 MPa, respectively. Initial σxx levelin two-pass weld 1 before treatment was lowerand was equal to 87 MPa. This is due to localtempering of weld metal deposited in the firstpass after making the second pass. After EDTσxx values did not exceed 6.5 MPa in the meas-ured zone that is comparable with the stress levelin the base metal before deposition. Changes ofσxx values in welds 1 and 2, depending on the

Figure 3. Change of values of stresses σxx in single-pass 1(a) and two-pass 2 (b) welds depending on the number ofcurrent discharges n

Table 1. Modes of EDT of welded joints of magnesium alloyML10 (capacitor storage C = 6600 μF, discharge ratio tr = 60 s)

EDTmode

number

Chargingvoltage U, V

Chargingcurrent* I,

A

Electrodepressure* P, N

Dischargetime* td, ms

1 200 1195 2792 1.2

2 500 3080 20461 1.6

*Procedure of determination of EDT parameters is described in [4].

30 11/2012

number of current pulses n are shown in Figure 3,from which it is seen that the maximum effec-tiveness of electrodynamic impact is achieved af-ter the first current discharge (n = 1) that allowslowering initial σxx values by more than 50 %.

Results of experiments conducted on AEICfragments lead to the conclusion that EDT of repairdeposits in the zone of characteristic damage of thestructure allows lowering the level of initial weld-ing stresses practically to base metal level.

EDT of full-scale AEIC was conducted in thelocations of repair cladding in the areas of postdamage (see Figure 1, b) and reinforcement forthe flange of cooling pipeline fastening (Fi-gure 4). EDT was performed in modes shown inTable 1 in the sequence corresponding to treat-ment of full-scale samples. During EDT cycleinitial stress level was recorded before and aftercladding, as well as after EDT. Analysis of cur-rent measurements of parameters of repair depositstressed state leads to the conclusion that afterEDT the level of stressed state in repair depositsis close to that of AEIC base metal.

It should be noted that manual tool for EDT(see Figure 4) enables access to AEIC repairwelds in all the positions. Power source for EDT,the weight of which does not exceed 3 kg, isquite compact, that allows placing it on the sur-face of the treated structure in the working zoneof EDT operator. EDT operators are exposed tothe impact of pulsed electromagnetic fields. Thisis related to the fact that the tool, which is thesource of magnetic radiation, is in direct contactwith the operator’s hand during EDT. Values ofintensity H of the magnetic field (MF) shouldnot exceed limit permissible levels (LPL) speci-fied by «State Sanitary Norms and Rules of Op-eration with Electromagnetic Field Sources»(DSN 3.3.6.096—2002). Determination of MF pa-rameters, corresponding to AEIC treatmentmodes, is an urgent task, related to taking meas-ures for industrial safety of EDT operators.

The main MF source is a flat inductor, whichis part of the working tool [4]. Amplitude valueof MF intensity at EDT operator workplace de-pends on pulse current, dimensions and shape ofdischarge circuit, as well as the distance betweenthe performer and field source. Such MF sourcesas discharge circuit and capacitor storage modulewere not considered, in view of small values ofmagnetic radiation.

Proceeding from analysis of amplitude-fre-quency characteristics of current pulses, appliedat EDT [4], conditions of MF radiation at EDTare at the lower limit of radio frequency range.This allowed isolating a frequency range from 1up to 10 kHz, in which it is necessary to determineMF level, corresponding to electrodynamic im-pacts with charging voltage of 200—500 V.

A flat inductor was a source of MF radiation,and the operator’s wrist located at the distanceof 70 mm from the inductor, was selected as thezone closest to MF source.

Plate with deposited bead from ML10 alloywas used for evaluation of MP parameters.

Intensity H of pulsed MF was determined us-ing instrumentation system GFI-1 (Hall sensor),the analog signal from which was recorded byTDS-1002 oscillograph with Fourier transforma-tion function. Certified sensor and oscillographensured measurement of the spectrum of MF in-

Figure 4. AEIC EDT in the zone of repair cladding ofreinforcement for the flange of fastening the post coolingpipeline: 1 – reinforcement for flange; 2 – manual toolfor EDT; 3 – power source for EDT

Figure 5. GFI-1 system for measurement of pulsed MF in-tensity at EDT: 1 – power source for EDT; 2 – Hallsensor; 3 – flat inductor; 4 – welded joint sample; 5 –MF intensity recorder

11/2012 31

tensity H from 8 up to 16,000 A/m. Amplitudevalues of pulsed current were recorded using Ro-gowski’s belt by a procedure described in [4].Three ranges of MF frequency were studied atcapacitor storage discharge, namely: 0—5, 50—1000 and 1000—10,000 Hz. Values of chargingvoltage of capacitor energy storage, at which MFintensity H was measured, were taken to be equalto 200 and 500 V, that ensures the charge energyof 300 and 800 J, respectively, and is close toEDT parameters, used at treatment of AEIC re-pair welds. Hall sensor was fastened on inductorouter surface in the zone of operator hand locationthat allowed studying the parameters of horizon-tal and vertical components H of magnetic fluxat EDT. Recording H values was conducted dur-ing an isolated discharge of capacitor storagethrough an inductor mounted on a welded jointsample (Figure 5).

Values of pulsed current I and vertical com-ponent of intensity H of pulsed MF at EDT withcharging voltage of 200 and 500 V are shown inFigure 6. It should be noted that the ratio ofintensity H values in the vertical and horizontalplanes is equal to 10/1 that allows ignoring thelatter at calculation of MF characteristics.

It is found that amplitude values I at U = 200and 500 V are equal to 1200 and 3000 A, respec-

tively, and the time of current running does notexceed 1.4 ms (see Figure 6, curves 1). AmplitudeH values at similar U values are equal to 10,000and 30,000 A/m, respectively, and time of MFimpact is equal to 2.2 ms (Figure 6, curves 2).It should be noted that comparison of curves 1and 2, reflecting the ratio of values of pulsedcurrent and MF intensity during the current dis-charge shows that at I attenuation to zero valuesresidual magnetic flux was recorded in the meas-ured zone, the period of action of which is equalto 0.75—0.90 ms. At the moment of I achievingzero values, intensity H of residual MF at U == 200 and 500 V was equal to 4000 and10,000 A/m, respectively. Presence of MF aftercurrent action in the discharge circuit is over, isattributable to residual magnetization of flat in-ductor, as well as running of attenuating currentin a disc from non-ferromagnetic material, incor-porated into the working tool.

Proceeding from the obtained data, calcula-tion-based estimate of relative energy load (REL)in the studied spectrum of MF frequencies wasperformed by the following procedure [5]:

REL = Hm

LPL, (1)

where Hm is the MF intensity, A/m (Hall sensorreadings); LPL are the data from standard DSN3.3.6.096—2002.

Time of operator working top was assigned asan 8-hour shift, which is equal to 28,800 s. Fullperiod of action td of pulsed MF, as shown inFigure 6, a, was equal to 2.2 ms for all the studiedvalues of charging voltage.

Admissible values of operator exposure tadmand number of tool switching on operations nadmin the studied MF were calculated by the follow-ing procedure [5]:

tadm = top

2∑REL, (2)

nadm = tadm

td. (3)

Data of calculation of MF parameters, givenin Table 2, lead to the conclusion that at chargingvoltage of up to 200 V, EDT operator can performnot more than 1100 actions of thermodynamicimpact per a work shift, and at 500 V voltage –not more than 100.

The number of electrodynamic impacts perone item does not exceed 20—30 discharges. Thus,production cycle of AEIC reconditioning, includ-ing EDT, provides safe working conditions forEDT operators under the condition of chargingthe capacitor storage up to maximum voltagevalue of 500 V.

Figure 6. Amplitude values of pulsed current I (1) andmagnetic field intensity H (2) at charging voltage of 200(a) and 500 (b) V

Table 2. Spectral composition and relative energy load of MF atAEIC EDT (discharge time td = 0.0022 s)

Chargingvoltage U,

V

MF RELAdmissibledischarge

time tadm, s

Admissibledischargenumber

nadm

Frequency range, Hz

0—5 50—1000 1000—10000

200 0.64 4197 1705 2.45 1100

500 8.35 51968 13426 0.22 100

32 11/2012

CONCLUSIONS

1. Technology of repair welding of damage inAEIC from magnesium alloy ML10 was devel-oped, including EDT of welds to lower the levelof residual welding stresses.

2. By the results of EDT of full-scale AEICfragments with characteristic damage of the itemreconditioned by repair welding, it is establishedthat EDT allows eliminating residual stresses inthe weld.

3. Experimental procedure was developed, onthe basis of which the influence of charging volt-age on magnetic field intensity at EDT of weldedjoints of magnesium alloy ML10 was studied.

4. It is established that at up to 200 V chargingvoltage EDT operator can perform not more than1100 actions of electrodynamic impact per a

working shift, and at the voltage of 500 V –not more than 100, that supports the productioncycle of repair welding of AEIC from magnesiumalloy ML10.

1. Stepanov, G.V., Babutsky, A.I. (2007) Modeling ofstress relaxation under action of pulsed electric currentof high density. Problemy Prochnosti, 2, 113—120.

2. Antonov, Yu.A., Ragozin, Yu.I. (2001) Pulse methodof residual stress relieving. Fizika i Khimiya Obrab.Materialov, 3, 91—95.

3. Lobanov, L.M., Pashchin, N.A., Loginov, V.P. et al.(2007) Change of the stress-strain state of welded jo-ints of aluminium alloy AMg6 after electrodynamictreatment. The Paton Welding J., 6, 7—14.

4. Lobanov, L.M., Pashchin, N.A., Cherkashin, A.V. etal. (2012) Efficiency of electrodynamic treatment ofaluminium alloy AMg6 and its welded joints. Ibid.,1, 2—6.

5. Levchenko, O.G. (2010) Occupational safety and he-alth in welding engineering. Kiev: Osnova.

PROCEDURE FOR CALCULATIONOF DIMENSIONS OF NOZZLES IN WELDING

USING TWO SEPARATE GAS JETS

V.M. BELOKON and A.O. KOROTEEVState Institution of Higher Professional Education «Belarusian-Russian University», Mogilyov, Republic of Belarus

Advantages of the process of welding without short-circuiting with double gas shielding of the arcing zoneare shown. The arc is shielded by argon, and the weld pool – by carbon dioxide gas, fed through twoconcentrically located nozzles. Calculation of arc radius in its largest cross-section was performed. Calcu-lation of weld pool length allows determination of the diameter of nozzle for carbon dioxide feed. Applicationof higher welding parameters requires increasing the diameter of nozzles, which can be calculated by similarprocedures.

Keywo r d s : arc welding, consumable electrode,shielding gases, separate jets, dimensions of nozzles, cal-culation procedure

Gas shielded welding finds wide application inproduction of various structures. At that CO2welding or welding in its mixtures with oxygen,argon etc. are often preferred. Welding withoutshort-circuiting with double gas shielding, i.e.welding arc is shielded by Ar and weld metal byCO2 is presented to be promising method. Thismethod allows significantly reducing losses forelectrode metal spattering, expenses for cleaningof near-weld zone from spatters and shielding gascosts [1—4].

Main parameters of each jet of shielding gaswere experimentally determined by a number ofdomestic and foreign researchers and recommen-dations were provided for selection of dimensionsof welding torch nozzles [5 et al.].

The aim of the present work is a developmentof procedure for calculation and determinationof dimensions of nozzles (for Ar and CO2) inreversed polarity current welding with two radialjets of shielding gases.

An electric arc consisting of three areas (an-ode, cathode and column) is used as a power sourcein consumable electrode welding. The anode andcathode areas have small dimensions. Anode spotin Ar welding can cover the whole end surface ofthe electrode and transfer to its side surface. Atthat transfer of electrode metal takes place in aform of small drops or jet that has positive effecton process of the electrode metal transfer, reducingspattering and splashing.

Argon shield of the cathode and anode areas,as well as arc column, can provide welding proc-ess, connected with positive effect of arcing inargon in welding with two concentric gas flows.

Putting of arc column to homogeneous chan-nel with uniformly distributed within it tempera-© V.M. BELOKON and A.O. KOROTEEV, 2012

11/2012 33

ture and current is a reasonable idealization beingsufficiently close to practice and does not violat-ing the main physical representations in series ofassumptions, considering that welding arc burnsin iron vapors (in steel welding). Figure 1 showsthe accepted scheme known as «channel» modelof arc column [6, 7]. According to this model anaverage current density in the arc is distributedalong the section with effective radius ref.

The average current density based on «chan-nel» model is determined on formula

jav = 5.5⋅10—8 Ui

38/12a2/3

ge2/3Ia

1/3, (1)

where Ui = 7.87 V is the ionization potential ofmetal vapors; a is the ratio of static weights ofions and atoms of iron vapors (a2 = 12/5); ge == 35⋅10—20 m2 is the section of collision of atomswith electrons in Ar welding; Ia is the weldingarc current, A.

Effective radius of arc column

ref = √⎯⎯Ia

πjav =

2.4⋅103Ia2/3ge

1/3

U19/12a1/3. (2)

The whole arc current according to «channel»model passes through section with radius R de-termined on formula

R = 2R0.5, (3)

where R0.5 is the conventional arc radius relatedwith effective radius by relationship

ref = 1.4R0.5. (4)

The next will be obtained solving simultane-ously expressions (2)—(4):

R = 3.4⋅103 Ia

2/3 ge1/3

U19/12a1/3 . (5)

The results of calculation of effective and ac-tual radius of the arc column depending on weld-ing current intensity, represented in Figure 2,show that an internal nozzle feeding argon jet ofaround 9 mm diameter (R = 4.5 mm) is suffi-cient for complete shielding of welding arc fromambient environment by argon using normalwelding modes (up to 450 A). The diameter ofnozzle should be increased for welding modesof ≥450 A.

Shielding of welding arc only is obviously notenough for obtaining of quality weld. Shield ofa surface of weld pool molten metal from inter-action with atmosphere is necessary to be pro-vided.

Figure 2. Dependence of effective ref and actual R radiusesof arc column on welding current

Table 1. Width of weld and active zone in welding using 1.2 mmdiameter wire

Weldingcurrent, A

Arc voltage,V

Weldwidth,mm

Active zonewidth, mm

Dimension ofactive zone 2R

acc. formula (5),mm

250 27 7.33 5.62 6.30

300 30 8.47 6.49 7.12

350 35 9.88 7.57 7.89

400 38 11.01 8.43 8.63

Table 2. Width of weld and active zone in welding using 1.6 mmdiameter wire

Weldingcurrent, A

Arc voltage,V

Weldwidth, mm

Active zonewidth, mm

Dimension ofactive zone 2Racc. formula

(5), mm

250 28 7.46 5.72 6.30

300 31 8.60 6.59 7.12

350 33 9.60 7.35 7.89

400 35 10.55 8.10 8.63

450 37 11.51 8.83 9.33

Figure 1. Scheme of «channel» model of arc column: ref –arc effective radius; jav – average density of arc current;T – average effective temperature of arc; R0.5 – conven-tional radius of arc column

34 11/2012

The weld pool consists of specific zones. Acentral zone includes a head part of the weldpool and part of a tail. Cross-section of the centralactive zone, based on some sources, coincideswith weld width [8]. In fact it is somewhatsmaller.

Weld width can be determined on the follow-ing formula [9, 10]:

b = 2√⎯⎯⎯⎯⎯2qπecγvwT

,(6)

where q = ηIaUa is the effective heat power ofarc; Ua is the arc voltage, V; η = 0.8 is theefficiency; cγ = 4.8 J/(cm3⋅K) is the volumetricheat capacity; vw is the welding speed, m/h; Tis the steel melting temperature, K.

Formula (6) can also be used for calculationof cross dimension of the weld pool active zone.For this temperature equal to metal evaporation

temperature should be assumed in this formula.Calculation results show that the dimension ofactive zone virtually coincides with the actualdimension of arc column section, calculated onformula (5), through which all the arc currentpasses. Deviation makes not more than 5—10 %.Tables 1 and 2 show the results of calculation ofthe weld width and cross dimension of the weldpool active zone.

Weld pool length is determined on formula [10]

L = q

2πλT, (7)

where λ = 47 W/(m⋅K) is the coefficient of heatconductivity of steel.

Radius of outer nozzle for CO2 feed, consid-ering shield of surface of the weld pool from

Figure 3. Dependence of dimensions of torch nozzles on welding current in welding using 1.2 (a) and 1.6 (b) mm diameterwire

Table 3. Results of calculation of weld pool length and radius ofouter nozzle in welding using 1.2 mm diameter wire

Weldingcurrent, A

Arc voltage, VLength of weld

pool, mm

Radius of outernozzle acc. formula

(8), mm

250 27 13.07 9.40

300 30 17.43 13.20

350 35 23.73 18.79

400 38 29.44 23.94

Table 4. Results of calculation of weld pool length and radius ofouter nozzle in welding using 1.6 mm diameter wire

Weldingcurrent, A

Arc voltage, VLength of weld

pool, mm

Radius of outernozzle acc. formula

(8), mm

250 28 13.56 9.82

300 31 18.01 13.71

350 33 22.37 17.58

400 35 27.12 21.84

450 37 32.25 26.50

11/2012 35

interaction with atmosphere, can be calculatedusing formula

Rout = L — b/2. (8)

Results of calculations are summarized in Ta-bles 3 and 4.

Formula calculations show that applicationof 20 mm radius nozzles is enough for arcs withup to 350 A welding current and 1.2 and 1.6 mmdiameter wires. Radius of the nozzle for outergas jet should be larger in the case of weldingwith higher currents. The outer nozzle for thepurpose of economy of shielding gas can be madein ellipse form, the cross dimension of whichequals the weld pool and the longitudinal dimen-tion equals its length.

The next sequence of calculation is proposedfor welding torch nozzles: calculation of averagecurrent density in the arc column; determinationof effective and actual radiuses of the weldingarc; after that using values of these parametersdetermination of diameter of nozzle for argonfeed (see Figure 2) and dimensions of the weldpool; determination of diameter of outer nozzlefor CO2 feed (Figure 3) considering the dimen-sions of active zone. Diameter of the outer nozzlecan be reduced considering spreading of shieldinggas flow in welding of flat joints [11].

Diagrams, shown in Figure 3, simplify theprocesses of practical fulfillment of the proposedprocedure.

Analysis and calculations performed alloweddetermining the optimum relationship of shield-ing gases which should make 1:4, i.e. 20 % Arand 80 % CO2, from the general necessary con-sumption.

Primitive technical and economic calculationsshow that welding with two separate gas jets canbe applied not only to special materials, but tolow-alloyed as well as low-carbon steels. Econ-omy only of electrode metal at that makes 20—95 kg per 1 t of wire that is character for pureargon welding without short-circuiting and itcovers an insignificant increase of shielding gas

consumption in comparison with CO2 welding,welding with double and triple mixtures. Fivetime reduction of argon consumption is observedin comparison with pure argon welding.

CONCLUSIONS

1. Sequence and procedure for calculation of di-ameter of welding torch nozzle in welding withtwo separate jets of shielding gas is proposed.The optimum ratio of Ar and CO2 in generalconsumption of shielding gas makes 1:4, i.e. Ar ++ 80 % CO2.

2. Calculation of diameter of nozzles for argonfeed based on «channel» arc model was carriedout, and calculation of dimensions of nozzles forCO2 feed was performed considering the weldpool dimensions. It was determined that internalnozzle feeding argon jet of around 9 mm diameteris enough for welding in normal modes (up to450 A).

1. Laevsky, V.S., Dyurgerov, N.G., Lenivkin, V.A. etal. (1969) Consumable electrode welding of low-carbon steels in combined shielding. Svarochn. Proiz-vodstvo, 10, 21—22.

2. Potapievsky, A.G. (1973) Consumable-electrode gas-shielded welding. Moscow: Mashinostroenie.

3. Paton, B.E. (1974) Technology of fusion arc weldingof metals and alloys. Moscow: Mashinostroenie.

4. Malyshev, B.D., Akulov, A.I., Alekseev, E.K. et al.(1989) Welding and cutting in industrial engineer-ing. Ed. by B.D. Malyshev. Vol. 1. Moscow: Stroj-izdat.

5. Fedko, V.T., Kiyanov, S.S., Shmatchenko, V.S. etal. (2003) Application of double-jet nozzles for gas-shielded welding. Avtomatizatsiya i Sovr. Tekhnolo-gii, 3, 12—18.

6. Leskov, G.I. (1971) Welding electric arc. Moscow:Mashinostroenie.

7. Pavlyuk, S.K., Belokon, V.M. (1974) About stabil-ity of arc ignition process in consumable electrodewelding. Svarochn. Proizvodstvo, 4, 51—53.

8. Erokhin, A.A. (1973) Principles of fusion welding.Moscow: Mashinostroenie.

9. Akulov, A.I., Belchuk, G.A., Demyantsevich, V.P.(1977) Technology and equipment for fusion weld-ing. Moscow: Mashinostroenie.

10. Petrov, G.A., Tumarev, A.S. (1967) Theory of weld-ing processes. Moscow: Vysshaya Shkola.

11. Popravka, D.L., Khvorostov, N.E., Proskurin, V.N.(1973) Some principles of flow of shielded gas jet.Svarochn. Proizvodstvo, 6, 33—36.

36 11/2012

BRAZE-WELDING WITH WELD METAL PEENINGDURING ITS SOLIDIFICATION

A.S. PISMENNY, I.V. PENTEGOV, V.M. KISLITSYN, E.P. STEMKOVSKY and D.A. SHEJKOVSKYE.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

Advantages of the process of braze-welding of zinc-plated steel, including impact peening of metal in thejoint zone at metal cooling stage, are considered in comparison with the widely accepted process of brazingby copper-based filler metals with arc heating.

Keywo r d s : resistance spot welding, braze-welding,peening, explosion compression mechanism, joint strength

At present the process of brazing by Cu-based fillermetals with arc heating became widely acceptedfor joining sheet Zn-plated steel. Selection of thisvariant of the joint is due, unfortunately, not tothe desire to achieve a high quality of the joint,but to a possibility of improvement of process ef-ficiency using currently available equipment forautomatic or semi-automatic arc welding in theatmosphere of active shielding gases [1].

Main difficulties in welding Zn-plated steelare caused by that zinc starts evaporating muchearlier than base metal melting temperature. Be-cause of appearance of zinc vapours over the weldpool, electric arc loses its stability that promotesappearance of weld porosity, undercuts and otherdefects.

In this connection, in arc welding of Zn-platedsteel the mode of heating with lower heat inputis used, and Cu-based alloys, for instance ofCuSn6P, CuSi3, CuSi2Mn, CuSi3 composition,are applied as filler wire [1, 2].

As in case of application of the above fillermaterials it is possible to avoid base metal melt-ing, joints of this kind can be included into thecategory of brazed joints, even though in thisprocess the fluxes, which are compulsory for con-ducting the process of brazing in an uncontrolledatmosphere, are not used. Here, the arc is a heatsource, comparable in its intensity with the heatevolved, for instance, at flame heating.

Unfortunately shielding gasses (argon, he-lium) used in arc welding do not provide thenecessary degree of wetting and spreading of Cu-based filler metal over Zn-plated steel surface.This circumstance leads to appearance of under-cuts in the joint zone, and promotes initiation ofmicrocracks lowering the joint strength at itsoperation under the conditions of cyclic load ap-plication.

In view of such shortcomings of the processof brazing Zn-plated steel with arc heating, thequestion of selection of an optimum variant comesup, which would provide not only high efficiencyof the technological process, but also high qualityof the joint.

One of the promising variants of producingjoints of coated metals is braze-welding, whichis a unique method to produce joints of similar-and dissimilar metallic and nonmetallic materi-als. A significant difference of braze-weldingfrom other joining methods is preliminary addi-tion of low-melting (compared to materials beingjoined) interlayer between the materials beingjoined or its formation during heating. In thecase of joining Zn-plated steel, such an interlayeris the zinc coating which melts at the temperaturemuch lower than that of steel melting.

In addition, braze-welding process is charac-terized by application of single or multiple com-pression force (peening), required for removal ofthe greater part of low-melting interlayer fromthe joint zone, that greatly increases jointstrength.

Attempts to apply metal peening at the finalstage of the welding process to improve weldedjoint strength were realized, for instance, in theunits for resistance spot welding of metal [3, 4].

This variant became applied in welding ofmetals prone to formation of cracks, loosenessand pores, in order to improve the fatiguestrength of welded joints. However, in practicethis kind of «peening» did not lead to any no-ticeable increase of joint strength, because of thelow speed of compression force application,caused by the use of a pneumatic drive of dis-placement of mobile welding electrode and iner-tia of its suspension assembly. As a result, insteadof high-speed peening the weld spot metal wasexposed to static compression force.

Delaying of the moment of peening force im-pact on the weld spot metal at temperature below

© A.S. PISMENNY, I.V. PENTEGOV, V.M. KISLITSYN, E.P. STEMKOVSKY and D.A. SHEJKOVSKY, 2012

11/2012 37

the points of structural transformations of treatedmetal crystalline lattice, for instance, point Ac1,did not lead to any significant deformation ofweld spot metal or development of mechanicalcompressive stresses in the near-weld zone. Inaddition, application of peening force to solidi-fied metal did not promote removal of the greaterpart of low-melting interlayer from the jointzone, particularly, in braze-welding processesthat might have essentially improved the jointstrength.

Improvement of the effectiveness of peeningof welded or brazed joint metal turned out to bepossible due to an abrupt increase of the speedof compressive force application. Results of theconducted comparative studies of the producedjoint strength show the obvious advantages ofimpact application of the compressive force, com-pared to the strength of joints, produced at ap-plication of static compressive force [5].

Impact peening of metal in the joint zone,conducted at its solidification stage, leads to de-velopment of several technologically importantphenomena, promoting an improvement ofwelded joint strength.

The features of this variant of thermomechani-cal treatment of metal in braze-welding or weld-ing processes include development of mechanicalcompressive stresses, both in the connecting weldmetal, and in the HAZ, which are preserved inthe process of further cooling of the metals beingjoined. In addition to that, improvement ofwelded joint strength is further promoted by theprocesses of refinement of metal crystalline struc-

ture in the joint zone caused by high-speed de-formation at higher temperature.

Results of technological studies described in[5] were obtained at application of an electro-magnetic drive of the system of compressive forceimpact. However, inertia of the mobile part (sus-pension) of welding electrode limited the speedof compressive force application on the level of300 m/s [6].

The assumption of a good potential for in-creasing the upsetting speed in the processes ofresistance spot welding or braze-welding is basedmainly on the experience of forge welding withindirect heating of the metals being joined. More-over, an abrupt increase of the compressive forcespeed, in all probability, should inevitably beaccompanied by appearance of new technologicaleffects.

Speed of compressive force application can beincreased by using an explosion of hydrogen-oxy-gen mixtures, in which the velocity of propaga-tion of the shock wave front reaches 3000 m/s[7]. Even higher speeds (up to 6000 m/s) canbe achieved at application of the electrohydrauliceffect [8].

This work presents a variant of the process ofresistance spot braze-welding with application ofa compressive force on weld metal, which is cre-ated as a result of an explosion of hydrogen-oxy-gen mixtures of a stoichiometric composition,produced by electrolysis of water in portable gasgenerators. For instance, a generator of P-105type which was developed by PWI for flame braz-ing and welding of small-sized products providesthe efficiency of hydrogen-oxygen mixture of upto 350 l/h at up to 0.07 MPa excess pressure.

Schematic of explosive drive of displacementof electrode assembly of the unit for spot braze-welding, which does not differ so much from theknown schematics of the units using peening, isgiven in Figure 1. As shown in the Figure, thehydrogen-oxygen mixture is generated in electro-lyzer 1 and comes to the electrode assemblythrough electropneumatic valve 2, which isswitched on by controller 3. The latter ensuresswitching on the heating current, regulation ofthe number of alternating current pulses, timeof delay of electropneumatic valve switching onafter completion of a series of current pulses forheating the samples being joined, switching ongenerator 4 of high-voltage pulses applied to thedevice for firing the combustible mixture 5. Elec-trode component is made in the form of cylinder6, the lower part of which accommodates bellows7 with the assembly for fastening replaceablewelding electrode 8 and drainage hole for dis-charging the combustion products 9.Figure 1. Schematic of the drive of electrode assembly for

spot braze-welding (for 1—9 see the text)

38 11/2012

The assembly of fastening the welding elec-trode is made on the basis of bellows, the corru-gations of which close completely at the initialmoment to ensure the preliminary compressiveforce, but allow additional displacement of weld-ing electrode in the vertical direction under theimpact of pressure pulse formed in the explosionchamber at firing of the combustible mixture.

Process cyclogram (Figure 2) shows the se-quence of impact of the following technologicalparameters: application of compressive force(preliminary P1 and impact action P2); curve oftemperature rise in the joint zone when goingthrough four half-cycles (1—4) of heating cur-rent; time interval of delay of the impact of com-pressive force t1 and moment of application ofhigh voltage to the firing device of the systemof electrode suspension with an explosion driveof electrode assembly of the unit for spot braze-welding (in this case a unit pulse of peening forceapplication is implemented) or with an electro-magnetic drive, providing the impact of severalpulses of the peening force.

As is seen from the cyclogram, after creatinga preliminary compressive force in the joint zone,the electropneumatic valve connecting the gasvolume of electrolyzer with the explosion cham-ber operates at switching on the heating current.Blowing of explosion chamber volume throughthe drainage hole is performed, with its sub-sequent filling with the combustible mixture. Af-ter disconnection of explosion chamber volumefrom the electrolyzer volume the controller en-sures connection of electric circuit of powersource of combustible mixture firing system.Drainage hole remains open after blowing of ex-plosion chamber volume, as it was experimentallyestablished that the small diameter of this hole(about 1 mm), high velocity of explosion wavepropagation and short time interval between themoment of completion of explosion chamber fill-

ing and moment of application of the high-volt-age pulse of mixture firing almost do not lowerthe effectiveness of the explosion wave.

Welding head with the mechanism of impacttreatment of the weld of explosion type is shownin Figure 3.

Figure 4 gives a typical oscillogram of arrange-ment of welding current pulses (four half-cyclesin this case) and pulses of current, ensuring firingof the combustible mixture.

Signal for firing the combustible mixturecomes from the controller of welding circuitpower unit after counting the time of delayingthe time interval from the moment of completionof welding current pulse, set by the operator.This completes the cycle of welding an individualspot and it is repeated many times in this sequenceat process realization, similar to the process ofresistance seam brazing.

Technological studies of the process of braze-welding with explosive application of unsettingforce were performed using samples of Zn-coatedSt3 steel 0.3 mm thick. Process of braze-weldingwas conducted in a laboratory unit for resistance

Figure 2. Cyclogram of the process of braze-welding withimpact application of the peening force

Figure 3. Welding head with explosion type mechanism ofimpact treatment of the weld

Figure 4. Oscillogram of welding current pulses and currentpulses in the circuit of explosive application of the com-pressive force

11/2012 39

spot welding, in which the upper electrode holdersupported the electrode assembly with an explo-sion upset drive.

Results of technological experiments show theability to preserve the continuity of zinc coatingin the zone of contact of the sheets being joinedwith welding electrodes due to lowering of thelevel of heat evolution on transient resistancesand intensive heat removal from this zone intoforcibly cooled electrodes, ensuring the compres-sive force in braze-welding zone.

Impact action of pressure at the stage of metalsolidification in the joint zone promotes anabrupt increase of heat removal from the surfacelayers of the metals being joined, which weresubjected to peening, thus leading to narrowingof the HAZ.

As is seen from the photograph of macrosectionof brazed joint with zinc coating (Figure 5), di-mensions of HAZ can be essentially reduced evenat joining sheets of greater thickness (0.4 mm).

One of the features of braze-welding processwith impact upsetting in some cases is an almostcomplete ousting of the low-temperature interlayerat minimum thinning of the joined sheets in thezone of Zn-plated sheet joint (Figures 5 and 6).

As shown by experimental results, the effec-tiveness of impact application of upsetting forcecan be realized not only at heating by electriccurrent, but also for other variants of joint zoneheating, for instance, flame, microplasma andarc. Advantages of flame heating (hydrogen-oxy-gen mixtures) include high accuracy of the pa-rameters, purity of gas mixture, and possibilityof fine adjustment of the required thermal en-ergy, applied to the joint zone (in the pulsedheating mode).

Possibility of application of indirect heatingis the only technological variant allowing weld-ing of nonmetallic materials to be performed.

In order to realize the variant of indirect heat-ing of the parts being welded it is sufficient tofit the welding head with a plasmatron or flametorch. In this case the controller should ensure

a continuous sequence of commands for feedinga thermal energy pulse of specified value to theheating zone and impact treatment of the jointzone. In the future transition from pulsed heatingto continuous heating is possible, which can becontrolled by selection of heating source power,speed of welding head displacement, and distancebetween indirect heating source and joint zone.

CONCLUSIONS

1. Impact peening of metal during spot braze-welding of Zn-coated sheet steel is not only themost acceptable, but also the only, in our opin-ion, variant of producing the joints, the goodprospects of which are proved by the possibilityof preservation of the initial coating layer afterheating of the sheets being joined above zincmelting point and improvement of corrosion re-sistance of the joints in service due to eliminationof copper and its alloys from the joint zone.

2. Impact application of the compressive forceduring welding at the stage of metal cooling inthe joint zone promotes an increase of braze-welded joint strength due to refinement of metalstructure, and lowering of the probability of de-fect development in the joint zone, particularly,pores, cracks, and gas inclusions.

1. (2010) Arc brazing of galvanized sheets. Weldingand Cutting, 9(1), 20.

2. (2010) Investigations into the material failure of re-sistance weld spots on ultrahigh-strength steels.Ibid., 9(3), 167—173.

3. (2000) Equipment for resistance welding: Refer.Book. Ed. by V.V. Smirnov. St.-Petersburg: Energoa-tomizdat.

4. Moravsky, V.E., Vorona, D.S. (1985) Technologyand equipment for spot and projection capacitor-dis-charge welding. Kiev: Naukova Dumka.

5. Pismenny, A.S., Kislitsyn, V.M. (2010) Influence ofweld metal impact treatment on welded jointstrength. The Paton Welding J., 1, 38—40.

6. Pismenny, A.S., Pentegov, I.V., Kislitsyn, V.M. etal. (2011) Devices for impact treatment of the weldin the process of resistance spot welding. Ibid., 1,45—47.

7. (1973) New technological processes in precise instru-ment engineering. Ed. by R. Zeviga. Moscow: Energiya.

8. Yutkin, L.A. (1955) Electrohydraulic effect. Lenin-grad: Mashgiz.

Figure 5. Typical macrosection (×50) of brazed joint withimpact peening of the joint zone

Figure 6. Macrosection (×50) of braze-welded joint acrossthe section near the joint center

40 11/2012

COMPARATIVE ANALYSIS OF ISO 18841:2005 STANDARDAND RF 26389—84 STANDARD ON EVALUATION

TO HOT CRACK RESISTANCE IN WELDING

B.F. YAKUSHINN.E. Bauman Moscow State Technical University, Moscow, Russian Federation

Differences between the EU standard on tests to hot crack resistance in welding and standard 26389—84in force in the Russian Federation are considered, and specific variants for their harmonization to quantitativeevaluation of the sensitivity of steels and welding consumables to hot cracking are suggested.

Keywo r d s : arc welding, hot cracks, brittlenesstemperature range, ductility margin, crack physicalsimulation, critical rate of deformation, technologicaland machine methods of testing, EU and RF standards

The first standard of such designation GOST26389—84 was developed at Welding Faculty ofthe N.E. Bauman Moscow State Technical Uni-versity with the assistance of author. Monographof Prof. N.N. Prokhorov and works of other re-searchers studying the problems of hot crack(HC) resistance, the results of which were widelydiscussed during two meetings on problem of HCformation in the welds, castings and ingots in1958 and 1962 [1], became a theoretical basis ofthis standard.

As a result, a theory of production strengthof metals during solidification in welding wasstated. Based on this theory, the HC appear inthe alloys under effect of welding stresses in abrittleness temperature range (BTR) as a resultof exhausting of ductility margin δBTR in a periodof solid-liquid state tBTR. Possibility of HC for-mation is determined by ratio of three main fac-tors: BTR, minimum ductility δmin of metal inthe BTR, and intensity of deformation rise in theBTR, depending on rigidity of welded structure.If accumulated deformation εi exceeds currentvalue δi(T)min (Figure 1) in the BTR limits theHC will appear.

Critical tension speed Vcr, equal δ/tBTR ratio,was taken as an index of weld metal HC resis-tance for specific welding mode, and at compari-son of modes [2] critical tension rate Bcr, equalδmin/BTR, at which HC formation is possible,was considered. These indices are to be deter-mined by means of increase of deformation rateof welded samples from studied alloys up to HCappearance.

There are two variants of determination of Vcrand Bcr in GOST 26388—84:

• technological methods, i.e. by means ofwelding of the samples under conditions of in-crease of rigidity (thickness, level of fastening,mode of welding up to HC formation);

• machine methods, i.e. by means of increaseof deformation intensity of weld being solidifiedusing testing machine.

Testing machines of three types, i.e. LTP1-4,LTP1-6 and MIS, allowing tension or bendingof small-dimension samples in welding with ad-justed speed up to HC appearance [1], were de-veloped at the Bauman MSTU for practical ap-plication. This provided a wide implementationof a testing technique at the E.O. Paton ElectricWelding Institute, main research institutes, atplants, as well as abroad [3—5]. Effect of standard26389—84 on the territory of the Russian Federa-tion was reinstated since 2000.

ISO 17641 standard, developed by EuropeanCommittee for Standardization (CEN) in col-laboration with Technical CommitteeISO/TC44 «Welding and related processes» in2005, consists of preface and two independentparts. The preface of ISO 17641-1:2004 describesthe methods of testing to HC resistance and areasof their application.

The first part of standard ISO 17641-2:2005describes in details a test procedure applyingwelding of butt and tee samples of natural rigid-ity, and methods for testing of the welded sam-ples with forced loading are characterized in thethird part of ISO 17641-3:2005.

Its configuration corresponds with that of theRF standard. General favorable reception ofstandard ISO 17641 is given in work [5].

However, insignificant selectivity in compara-tive testing as well as inapplicability for testingof sheet samples of not less than 10 mm thicknessshould be noted proceeding from proposed pro-duction tests and types of welded joints of «natu-ral» rigidity. Therefore, formation of HC is un-© B.F. YAKUSHIN, 2012

11/2012 41

likely in testing of modern quality electrodes ofmany grades and modes of testing do not repro-duce welding conditions of more rigid structuresto significant extent. Besides, limitation of di-mensions on length of the tee joint with double-sided weld and gussets prevent performance ofautomatic welding and weld failure after weldingfor detection of cracks on fracture type. The proc-ess of result evaluation is extremely complicatedand long due to manufacture and testing of spe-cial rod and plate samples from the welded joints.No data on thickness of metal to be welded andmethod of fastening, which eliminats deforma-tion in welding, are provided by procedure ofapplication of the sample with butt weld.

In contrast, the RF standard provides for pos-sibility of wide variation of thickness of samplesbeing welded, modes and methods of weldingwhich make more effect on crack formation proc-ess in comparison with chemical compositionchange. The universal samples with butt and cir-cumferential welds in sheet plane are as wellrepresented for this in the RF standard togetherwith rigid single-sided tee weld [6]. They allowchanging a metal thickness (1—12 mm), diameterof circumferential weld, modes and methods ofwelding in a wide range and, thus, obtainingcritical values at which formation of HC in sam-ple welding is possible.

Change of width of plates to be welded in thesample with butt joint allows increasing high-temperature welding deformations up to the level

sufficient for obtaining of quantitative result oftesting from weld metal of any composition(GOST 26389—84) that is very important for con-sumable selection.

The second part of standard ISO/TR 17641-3is represented in a form of engineering report ontests with forced «loading» in welding and canbe considered as its first project. It contains de-scription of American procedures Varestraint,TransVarestraint, Gleeble, as well as PVR pro-cedure developed in Austria [7].

A series of notes should be made on this docu-ment.

1. High-speed deformation by bending of so-lidifying weld metal on Varestraint and Trans-Varestraint procedures violates the principles ofphysical simulation in sample testing and condi-tions promoting failure in real welded structures.This note also refers to Gleeble procedure, inaccordance to which speed of high-temperaturetension of investigated samples in the BTR makes0.15—0.25 m/s (6—10 inch/s) [7].

2. Evaluation of degree of deformation in man-drel bending on formula ε = h/2R is suitable forhomogeneous, i.e. isotropic material. However,solidifying metal in welding has double-phasestructure and defomations are accumulated alongthe grain boundaries which is a reason for HCformation.

3. Evaluation to HC resistance in bendingtesting of the welded joint from thin-sheet metal,including pass, is impossible since the criticalvalues are not achieved in mandrel bending.

4. Proposed criterion of total length of cracksLtot does not consider ductility margin of metalin it being the main factor of crack formation.

5. Diffusion processes and high-temperaturecreep preparing conditions for HC nucleation arelimited by a dynamic deformation of weld beingsolidified. Elimination of these processes in thedynamic deformation develops seeming increasedHC resistance that can result in unpredicted fail-ure of the welded structures.

It should be noted that as a rough approxi-mation HC length can only characterize the BTRvalue in the dynamic deformation. Another fac-tor, i.e. ductility margin δBTR, can not be evalu-ated by number of cracks and their length, there-fore, Ltot does not considered to be a qualitativecriterion of tendency to HC (see Figure 1).

Inter-grain ductility of metal in the BTR ac-cording to Gleeble procedure is proposed to beevaluated on degree of its change outside theBTR, i.e. in area of high-ductility weld metalcondition that violate validity of tests [8].

Figure 1. Scheme of weld and near-weld metal tests to HCresistance in welding by means of deformation growing de-veloped using test machine: Tw(t) – welding thermal cycle;Ts(t) – simulation of thermal cycle; V1—V3 – growingof deformation in the BTR; Vcr – deformation resultingin crack formation; δ(t) – predicted character of ductilitychange in the BTR; Vcr = εcr/Δt – critical rate of defor-mation

42 11/2012

The advantage of the RF standard is that Vcr == δ/tBTR index can be evaluated using a methodrequiring no direct measurement of ductility ofthe weld metal and the BTR limits. At that, thesample is subjected to continuous, i.e. static, de-formation in the temperature range from upperto lower limits of the BTR up to 0.5Tmelt tem-perature in the investigated section with weldbeing solidified.

At that, intergranular shifts are not accumu-lated in deformation out of the BTR. They appearonly in the BTR that is the main advantage ofstatic tension or bending in HC resistance tests.

Inclusion of procedure of static deformation(programmable deformation cracking PVR) pro-cedure in project of ISO/TR 17641-3 standardis its positive moment. However, testing of PVRsample having deposited bead of large length re-sults in heating and increase of length of thesample between the machine grips that distortpredicted linear distribution of deformationsalong the sample.

Besides, a local concentration of deformationsunder the arc, to large extent machine vm, isinevitable as a result local arc heating of thesample and reduction of metal resistance to de-formation. Its level depends on thermophysicalproperties of steels and alloys being comparedand duration of deformation can be smaller thanthe BTR time.

In the RF standard the similar tests to HCresistance, oriented along the weld axis («pal-ing»), are carried out with growing tension speedwithin one series of the samples. This allows de-termining Vcr index and using it in selection ofalloys and consumables.

Large number of mutually exclusive indices(Table) and absence of correlation coefficientsbetween them should be noted at general evalu-ation of ISO 17641 standard. This provides a

necessity of purchase and operation of large num-ber of testing machines.

One conditional index of HC resistance Vcr(mm/s) is regulated in RF standard. Compari-son of its values is possible at equal speed wBTRof metal cooling in the BTR. In other cases physi-cal index Bcr (mm/°C) equal Vcr/wBTR is de-termined. This index allows evaluating resistanceof weld metal and near-weld zone to formationof longitudinal and transverse cracks in differentwelding methods [8].

New model of testing machine MIS (Figure 2)equipped with a fixture for static tension andbending of samples (Figure 3) in process of weld-ing, welding head with possibility of movementalong x, y and xy axes (circum—ferential weld)and fixture for electric contact heating and ten-sion of the samples for evaluation of metal ten-dency to formation of HC in the near-weld zone

Figure 2. General view of machine MIS for HC testing onthe RF standard: 1 – box for control of test parameters,their imaging and registration; 2 – manipulator of weldinghead for its movement along x, y and xy axes; 3 – forcemeasure device; 4 – machine for welding of samples andtheir bending or tensile tests; 5 – machine for simulationof welding cycle in the samples and HC tests at coolingstage

Tests to HC, types of cracking and designation according to ISO 17641:2005 standard

Test type Type of cracking Results Designation

Method of deformation along the weld axis Solidification Ltot Base metal (selection and confirmation)Consumable (selection and confirmation)Welding procedure

Liquation Ltot

Ductility reduction Ltot

Method of deformation across the weld axis Solidification Ltot Selection of consumable. Welding procedure

Tensile test of flat sample along the weld(PRV test)

Solidification Vcr Metal selection. Multipass welded joints.Welding procedure. Metal combination

Liquation Vcr

Ductility reduction Vcr Selection and confirmation of material

Tensile test in hot state (GleebleTM) Solidification BTR

Liquation BTR

11/2012 43

was designed for performance of metal tests toHC resistance. Presence of dynamometer inpower mains and dilatometer allows applyingMIS machines for testing of joints to cold crackresistance etc.

CONCLUSION

1. Theory of production strength is the basis ofstandard for quantitative evaluation of metal toHC resistance. In accordance to it the HCs arethe result of exhaust of ductility margin in theBTR under effect of welding stresses and defor-mations.

2. Critical rate of deformation, determinedbased on a fact of exhaust of ductility of thesamples with weld in the BTR at static machinedeformation, is quantitative index of metal HCresistance.

3. Machine methods of evaluation of metal toHC resistance in testing of small dimension sam-ples should provide the possibility of physicalmodelling of conditions resulting in HC forma-tion at manufacture of real welded structures.

4. Machine tests on ISO/TR 17641-3 stand-ard project using impact bending of the sampleswith weld (Varestraint and TransVarestraint) orimpact rupture of the samples (Gleeble) do nothave sufficient ground, since convective and dif-fusion processes, determining metal ductility in

the BTR, are not considered in determination ofLtot and high-temperature ductility dip (BTR)and not suitable for qualitative evaluation.

5. Technological tests of metal to HC resis-tance on GOST 26389—84 provide for applicationof samples with butt and tee welds as well aswidely used [8] samples with circumferentialweld of various thickness (1—20 mm) that guar-antee their suitability for comparison of consu-mables and technological variants of weldingand, thus, widening their versatility in compari-son with samples of ISO 17641—2 standard.

Figure 3. Types of samples applied for MIS machine testing

1. (1991) Welding and welded materials: Refer. Book.Vol. 1: Weldability of materials. Ed. by E.L. Maka-rov. Moscow: Metallurgiya.

2. Yakushin, B.F. (1969) Assessment of technologicalstrength depending on welding conditions. Svarochn.Proizvodstvo, 1, 19—21.

3. Bernasovsky, P. (2005) Contribution to HAZ liquati-on cracking of austenitic stainless steels. Phenomenain welds. Berlin: Springer.

4. Zhelev, A. (1988) Complex thermokinetic approachin assessment of hot cracks during welding: Syn. ofThesis for Dr. of Techn. Sci. Degree. Sofia.

5. Wilken, K. (1999) Investigation to compare hot crac-king tests. IIW Doc. IX-1945—99.

6. Tsarkov, V.A., Chuprak, A.I. (2010) Ways of har-monisation of national and European standards forassessment of metal resistance to hot cracking.Svarka i Diagnostika, 6, 58—62.

7. Yakushin, B.F. (2003) Harmonisation of RF and EUstandards on weldability tests. Svarochn. Proizvod-stvo, 1, 39—43.

8. Steklov, O.I. (1976) Strength of welded structuresin aggressive media. Moscow: Mashinostroenie.

44 11/2012

MULTICHANNEL MICROPROCESSOR CONTROLLERFOR DATA COLLECTION FROM THERMOCOUPLES

V.V. DOLINENKO, T.G. SKUBA, O.Yu. VASHCHENKO and N.F. LUTSENKOE.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

Development of 9-channel microprocessor controller of the system of welding thermal cycle recording withdigital interface for communication with PC of Ethernet-100Base-TX is presented. CAT and WRT typethermocouples are used as temperature sensors. Application of Internet-protocol of data exchange with PCof TELNET type allows designing automated system for research performance in the field of multipasswelding at minimum cost.

Keywo r d s : multipass arc welding, thermocouple,microprocessor controller, Internet, TELNET, sigma-delta ADC

Performance of experimental studies of mechani-cal properties of welds in critical structures in-volves application of expensive welding equip-ment and consumption of considerable materialresources, such as welded samples metal, elec-trode wire, shielding gas, as well as power. There-fore, one of the main objectives at experimentperformance is ensuring reliable recording ofmaximum possible scope of information to obtainthe most complete idea of the nature of structuralchanges in the HAZ metal of the item being

welded. Of considerable interest is the informa-tion on dynamics of variation of spatial tempera-ture field in the welded item during performanceof mutlipass arc welding. The most widely ac-cepted in welding are contact methods of tem-perature measurement using chromel-alumel(CAT) and tungsten-rhenium (WRT) (VR5/VR20) thermocouples which allow measuringtemperature in the item up to 1300 (CAT) and2500 °C (WRT), respectively, with ±1 °C error.

At present use of corporate Internet networkis an effective method of creating computerizedsystems for research performance. During appliedresearch in the field of arc welding using Internet

Figure 1. Block-diagram of MCRTC: 1 – analog module; 2 – digital module; PU – power unit; GD – galvanicdecoupling; LCI – liquid-crystal indicator; KN1, KN2 – control buttons; FP – front control panel; PA – preamplifier;LED – light-emitting diode; TC – thermocouples; CJTS – sensor of «cold junction» temperature

© V.V. DOLINENKO, T.G. SKUBA, O.Yu. VASHCHENKO and N.F. LUTSENKO, 2012

11/2012 45

technologies it is possible to solve the problemsof both galvanic decoupling and reliable dataexchange between the controllers and PC. Eth-ernet-100Base-TX bus is usually used as physicalinterface, and Internet-protocols of TCP/IP andTELNET type ensure reliable connection and er-ror-free data transmission. In connection withthe fact that batch-produced thermocouple con-trollers with Internet connection are not avail-able now, PWI developed a microprocessor con-troller, designed for recording thermal cycles ofmultipass welding [1] (further on referred to asMCRTC) which ensures digitizing signals from9 thermocouples and issuing the results to PC byTELNET protocol with real-time labels (Fi-gure 1).

MCRTC is designed as an analog and digitalmodules, connected to each other through gal-vanically decoupled interface of SPI type. Digitalmodule uses 32-bit microprocessor LPC2378(NXP Company). Analog module is realized onthe basis of OA microcircuits of AD822 type and

24-bit sigma-delta ADC of AD7794 type (AnalogDevices Company). «Cold junction» tempera-ture is measured by thermistor B57861S with10 kOhm resistance (EPCOS Company). Rangeof measured temperatures for thermocouples ofCAT type is equal to 0—1300 °C, and for ther-mocouples of VR5/VR20 type it is 0—2500 °C.Sampling rate is 3 Hz for all the channels.

Application of galvanically decoupled inter-face of Ethernet-100Base-TX type allows placingMCRTC in the immediate vicinity of the objectof studies, and removing the PC to up to 100 mdistance, that will shorten the length of thermo-couple wires to 1 m, and minimize the level ofelectromagnetic noise and interference.

Information exchange with the computer isperformed through network protocol of TELNETtype, which ensures storing the data receivedfrom the controller into the file specified by theuser. MCRTC operation is controlled using threespecial commands: TIME – current time setting;TP – thermocouple type selection, andMEAS – sending thermocouple readings to thecomputer.

Calculation of thermocouple readings (in Cel-sius degrees) is performed by the known methodbased on the read value of thermocouple emfusing calibration charts [2] and allowing for coldjunction temperature. To assess the controllermetrological characteristics, constant voltagewas applied to its inputs through a resistancedivider from a galvanic power source. Voltage atcontroller input was monitored by digital volt-meter MASTECH MS8218 (measurement pitchof 0.001 mV). After program displacement com-pensation at shorted inputs, metrological char-acteristics of all the measurement channels prac-tically coincided. Differences of obtained «volt-

Figure 2. MCRTC metrological characteristics: a – mode of thermocouple of CAT type; b – mode of thermocouple ofWRT type (A1)

Figure 3. Recorded thermal cycle of multipass welding(fragment): ΔtAc3 = 15 s; t8/5 = 37 s

46 11/2012

age-temperature» characteristics (Figure 2) fromgraduation charts [2] did not go beyond the limitsof error ±0.2 °C.

Figure 3 shows a fragment of thermal cycleobtained in an experiment with multipass weld-ing of V-shaped groove by short sections, usingthe developed controller and subsequent dataprocessing by MS Excel program. A thermocou-ple of CAT type with 0.6 mm wire diameter wasused in the experiment, which was caulked in onthe surface on an item from low-carbon steel at5 mm distance from the groove edge. From Fi-gure 3 it is not difficult to assess the time ofmetal staying above point Ac3 (15 s), and timet8/5 of cooling from the temperature of 800 to500 °C (37 s). Maximum temperature in the con-trolled zone was 1240 °C.

Developed microprocessor controller used foridentification of mathematical model of the heatsource in MIG/MAG welding, as well as duringinvestigations of structural transformations insteel in multipass welding, demonstrated a highstability to welding interference and reliabilityof data transmission in PC.

NEWS

DEPOSITION OF NICKEL COATINGS ON COPPER PLATESOF MOLDS OF MCCB USING FRICTION STIR

SURFACING METHOD

Intensive wear-out of the internal surfaces of the moldstakes place in a process of continuous casting of steelthat makes extremely high demands to quality of sur-faces of the copper plates. Copper plates with thenickel coating have the highest working capacity. Thisresults in increased wear resistance of the copper platesin 3—4 times.

Technology for deposition of nickel on the copperplates of machines for continuous casting of billetsusing friction stir surfacing (FSS) based on frictionstir welding method was developed at the E.O. PatonElectric Welding Institute of the NAS of Ukraine.

Welding is performed by a face of rotating toolhaving an extended pin which moves in a weld metal

in welding direction. Plastification of the metal takesplace in the metal-to-tool friction along the butt ofsurfaces being welded that results in its stir and for-mation of weld.

Material of the surfacing tool should be high-temperature and heat-resistant that allows workingat 1000—1200 °C temperatures. Shape of the toolhas an important role at that. Thus, the best resultswere obtained in application of a tool with conicalpin.

Working tools were manufactured from ultra-hardmaterials, i.e. tungsten-cobalt hard alloys with mi-croadditions of refractory compounds and cubic boronnitride, and they had complex configuration.

Lap slot weld is formed in a tool movement. Suc-cessive overwelding of such welds with overlayingallows nickel surfacing on the copper plate.

1. Dolinenko, V.V., Skuba, T.G., Vashchenko, O.Yu.(2012) Microprocessor controller for registration ofthermal cycles of multipass arc welding. In: Abstr. of11th Int. Conf. on Instrument Engineering: State-of-the-Art and Prospects (24—25 April 2012, Kiev).Kiev: NTUU KPI, 223.

2. GOST R 8.585—2001: Nominal static characteristics.Instead of GOST R 50431—92, MI 2559—99. Introd.21.11.2001.

Process of coating using FSS method

Fragment of 10 mm copper plate deposited by 3.5 mmthick nickel

11/2012 47

TECHNOLOGICAL INNOVATIONS –BASIS FOR INCREASE OF COMPETITIVENESS

OF THE U.S. WELDING PRODUCTION

O.K. MAKOVETSKAYAE.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

The tasks and problems of materials joining in industrial production are given. The model of developmentand implementation into production of technological innovations, offered by the Edison Welding Institute(USA), was considered.

Keywo r d s : welding production, science, innova-tions, industrial consortium

The risk of loosing the leadership in the worldeconomy causes ever more anxiety among thestate and business spheres of the USA. In therecent years the country has lost its position inthe world rating of competitiveness havingmoved from the first place in 2009 to the fifthone in 2011, and in 2010 the USA let China havethe first place as to the volume of industrial pro-duction [1].

The sector of industrial production is the cor-ner-stone for the US economy. Its volumeamounts 11 % in the GDP of the country, andin the total volume of export the volume of goodsof industrial production exceeds 60 %. About 13.4Million people are engaged in industry whichamounts about 9 % of all employees. The laborpayment in the sector of industrial production is20 % higher than that in the other non-industrialsectors of economy.

Since 2008 the crisis remains the main problemfor the US economy. However negative tenden-cies in the economy of the country were revealedas long ago as 2001. By that time in the periodof one year more than 2.5 Million working placeswere reduced in the sector of industrial produc-tion. Among the most anxious tendencies in theUS sector of industrial production, experts out-line the following:

• reduction of output of industrial products.The volume of industrial sector in GDP of thecountry for the period of 2000—2010 was reducedfrom 17 to 11 %;

• reduction of number of working places. Inthe period of 2000—2010, 37 % (6.5 mln) of work-ing places were reduced in industry;

• reduction of volume of foreign trade (thevolume of USA at the world market decreased

from 19 to 11 % (2000—2010) that resulted ingrowing of the trade deficit;

• increase in prices on goods of industrial pro-duction (increase of costs connected with safetyand protection of environment, taxation, laborpayment, reclamation, etc. was reflected on theprice of ready products which became one of thefactors deteriorating its competitiveness at theworld market);

• lack of qualified personnel. Only in the fieldof welding production the lack of qualified weld-ers amounts 500 persons per year [2].

The technologies for materials joining are theindispensable part of industrial sector of econ-omy. Welding and related technologies of joiningare closely integrated into the production processof fundamental branches of industry and consid-ered as the key non-alternative technologies forthem. Taking into account such a decisive roleof joining technologies for the economy, theEdison Welding Institute (EWI) together withthe American Welding Society in 2010 initiatedthe wide-scale study of state-of-the-art and pos-sible ways for growing competitiveness of indus-trial production where the materials joining iswidely applied. In the frames of the project «Thefuture of materials joining in the North America»the survey of goods manufacturers of six leadingbranches of industry was conducted to determinethe basic problems of those branches and theirneeds in technologies for joining materials. Theresults of research were studied in 2011 at thefinal conference «The growth of competitivenessof industry: the future of materials joining in theNorth America» in the work of which the scien-tific, governmental and social organizations,leading manufacturers of welding technologies«Lincoln Electric», «Trumpf», «Miller Elec-tric», etc. took part. In the final document thebasic problems of materials joining in industrial© O.K. MAKOVETSKAYA, 2012

48 11/2012

production and tasks of its development in thenearest five years were defined.

Under the modern conditions of globalizationof world economy the only possible way for thegrowth of competitiveness in industry is innova-tive development, i.e. improvement of level ofdevelopments, reduction of time for implemen-tation of technical innovations into production,improvement of level of education and qualifica-tion of personnel. The innovative developmentof economy implies also a high level of correlationbetween science, production and personnel train-ing. According to the survey of enterprises of sixleading branches of US industry carried out bythe EWI, the problems available in the branchof materials joining are closely connected withthe solution of namely these tasks. The resultsof survey are given in Tables 1 and 2 [3].

According to the results of research givenabove over the recent years the volumes of ap-plication of new progressive structural materialsand their combinations will grow in all branchesof industry (see Table 1). This is the basic taskof the branch of automobile industry and powerengineering, it is also included into the specifiedfour tasks of other branches of industry. The de-signers and manufacturers are ever more inter-ested in the application of new materials whichimprove technical characteristics of products andreduce their cost. For example, the need in re-

duction of mass of automobile resulted in increaseof application of high-strength steels, aluminium,magnesium alloys and composites. The growthof application of new structural materials re-quires development of the new technologies forjoining (see Table 2). It was mentioned by rep-resentatives of all the surveyed branches of in-dustry, and for space-aircraft and military indus-try this problem is the most challenged. Accord-ing to estimates of correspondents it is also nec-essary to reduce the cycle period «R&D – im-plementation of new developments into indus-trial production», to find the ways of reductionof costs on development and implementation ofinnovations, creation of on-line system for in-forming about innovative developments in thefield of materials joining. In total, these tasksreflect the need in development of strategy fordevelopment of joining technologies (Table 2).

The next important task is involvement ofqualified personnel into the field of joining tech-nologies. According to the data of the US Sta-tistics Bureau the number of employees and spe-cialists of all welding professions in the periodof 2002—2009 decreased from 1,076,498 to968,037 people, or by 10.08 %. Nowadays thedeficit of welders on long-term contracts amountsapproximately 500 people per year. However thisnumber can be higher as far as mastering theprofessional skills of welder is required else in

Table 1. Problems and priority tasks in the field of materials joining in the USA for the nearest five years (four first ranks onbranches of industry are mentioned)

Problems and tasks

Rank on branches of industry

Automobileindustry

Oil-and-gasindustry

Militaryindustry

Aerospaceindustry

Heavy-machinebuilding

Powerengineering

Deficit of qualified engineers and specialists in thefield of quality control of joints

1 4

Deficit of qualified workers-welders and workers ofother professions

3 1

Growth of competitiveness in countries with lowpayment of labor

3

Increase of expenses for development andimplementation of new processes, products, methods

2 2

Increase of time for evaluation of joints quality 3

Broadening of application of new materials andtheir combinations

1 4 3 4 4 1

Implementation of new technological processes 2 1

Decrease of time from designing to puttingtechnologies into production

1 3 4

Creation of on-line system for informing aboutinnovative technologies and methods, providingaccess to them

2 3

Increase of requirements to the quality of jointsperformance

4 2 2

11/2012 49

more than 25 professions. According to the dataof survey confirmed by the statistic data, thebranches of US industry lack in qualified work-ers-welders, engineers and also other specialistsin the field of welding and quality control. Thus,the main problem in oil industry is the lack ofqualified engineers and specialists in the field ofquality control of joints, and in heavy-machinebuilding the main problem is the deficit of work-ers-welders (see Table 1). The lack of qualifiedpersonnel is closely connected with modern-ization of system of personnel training, develop-ment and implementation of system of constantimprovement of qualification for the specialistsof all professions [4].

The main source of innovations is R&D. Ac-cording to the estimates of American experts thetotal costs on financing of research and develop-ments in the world will grow by 5.2 % in 2012and reach 1.4 Trillion USD, among which thevolume of USA will amount 36 % or 436 BillionUSD. The industry finances 64 %, the federalgovernment 29 %, meantime 71 % of all R&D,carried out in the country, are performed in theindustry. Table 3 gives data on structure of dis-tribution of financing of R&D in USA on basicsources of financing and performers.

The field of R&D becomes ever more openedfor cooperation in USA as well as in the wholeworld. The data of Table 3 show considerablegrowth of finances by the industry of both itsown R&D as well as fundamental ones carriedout by academic organizations in the interests ofthe industry. The federal government attractsalso considerable investments into the industrialR&D and other organizations. According to thedata of survey of «R&D Magazine» 80 % of in-dustrial companies finance joint research withacademic organizations and other companies. Notonly industry but also federal government showsever growing interest in obtaining income fromthe investments into the R&D. If several yearsago only 10 % of companies planned and obtainedincome from investments into R&D, nowadaysalready more than 50 % of companies considerthis value as a key indicator of their activity.

The Act of Bayh-Dole accepted in 1980 laidgrounds for the new state scientific and technicalpolicy of USA directed to the growth of com-petitiveness of the national economy. The Actallowed transferring intellectual property, cre-ated at the federal costs, to such non-federal per-formers of R&D as universities, private compa-nies and other subjects, and also allowed exclu-

Table 2. Necessary industrial technologies of materials joining and other types of works (four first ranks on branches of industry arementioned)

Required technologies/types of works

Rank on branches of industry

Automobileindustry

Oil-and-gasindustry

Militaryindustry

Aerospaceindustry

Heavy-machinebuilding

Powerengineering

Development of technology for joining the newprogressive materials

1 3 1 1 3 3

Increase in number and improvement of level ofeducation of engineers and designers in the fieldof joining technologies

2 2 1

Development of arc welding process (efficiency,quality, etc.)

1 3 1

Development of new methods for joiningdissimilar materials

2 2 2

Providing the on-line access to the databases ontechnologies for materials joining

4

Development of more sensitive, accurate, reliablemethods of non-destructive testing

3 4 4

Development of high-efficient technologies forwelding thick-sheet materials

2

Improvement of methods for education of welders(making them more perfect, purposeful, lessexpensive)

4

Development of strategy of development of newjoining processes

4

Improvement of resistance welding technology(quality, reliability, etc.)

4

Development of additive industrial technologies 3

50 11/2012

sive licensing of inventions which is the key con-dition for their commercialization. This and an-other accepted acts and decrees of the govern-ment, state programs of USA and stimulated in-tegration of fundamental and applied science,strengthened interest of industry in performanceof fundamental research, activated the inter-dis-cipline research, changed the policy as appliedto research infrastructure [5, 6].

To stimulate the carrying out of technologicalR&D in the field of joining technologies, tostrengthen correlation among scientific and in-dustrial sector, to reduce significantly the timeand broaden the branches of implementation ofinnovation products, the EWI together with theInstitute of Industry (USA) created a model ofdevelopment and implementation of technologi-cal innovations into industry in the field of join-ing technologies and successfully approbated itin practice. The proposed model is based on theidea of creation of new organizational structureswhich could promote the closer integration of allparticipants of innovation process: from the ideato development commercialization and wide im-plementation of innovations into production,namely:

• focused industry consortia, and• manufacturing technology application cen-

ters.Consortium is the temporary union of indus-

trial enterprises interested in development of newprogressive technology. The members of consor-tium define the basic technological problemswhich require attention, form project programand team of performer. As the performers forsolution of different specific tasks the consortiumcan attract centers of development and imple-

mentation of industrial technologies, researchlaboratories, commercial structures and other or-ganizations. The support of innovations develop-ment to the stage of commercialization is per-formed by the state through the state programs.The implementation of innovations into industryimplies wide attraction of funds of industrialfoundations and other sources. Table 4 showsscheme of interaction of consortium and devel-opment centers and implementation of industrialtechnologies. This scheme demonstrates one ofthe possible ideas of functioning of consortium,i.e. the possibility of attraction in the course ofdevelopment of innovation for the solution ofdefinite tasks of specialized centers of develop-ment and implementation of industrial technolo-gies and material resources which have availablehighly qualified personnel and necessary materialresources.

The aim of the model of consortium developedby the EWI is to reveal the needs in new tech-nologies of materials joining arising in thebranches of industry, to realize the developmentof these technologies and develop programs ofpartner cooperation for the creation and wideimplementation of new technologies into produc-tion. The example of realization of this model inpractice is the Consortium of Additive Technolo-gies and the Consortium of Technologies of Nu-clear Power Engineering created by the EWI in2010.

For example, the Consortium of additive tech-nologies combined the efforts of large corpora-tions of US space-aircraft industry, the clientsof the EWI and other private, social and stateorganizations interested in the development andwide implementation into production of ad-

Table 3. Structure of distribution of financing of R&D in the USA on the basic sources of financing in 2012, Million USD (percentof changes by 2011)

Financing source

Performer of R&D

Federalgovernment

Governmental funds,centers, national

laboratoriesIndustry

National fund andother academicorganizations

Non-profitable organizations

In total

Federal government 29,152(—2.5)

14,666(—3.69)

37,577(—2.42)

37,440(093)

6817(—2.29)

125,652(—1.61)

Industry — 202(2.20)

237,487(3.37)

3868(26.49)

2129(8.89)

279,685(3.75)

National scientific fund and otheracademic organizations

— — — 12,318(2.85)

— 12,318(2.85)

Other governmental organizations — — — 3817(2.72)

— 3817(2.72)

Non-profitable organizations — — — 3491(2.70)

11,055(2.70)

14,546(2.70)

In total 29,152(—2.51)

14,868(—2.36)

311,063(2.63)

60,934(2.85)

20,001(1.55)

436,018(2.07)

11/2012 51

vanced additive technologies. In total the Con-sortium united 24 industrial members and part-ners for performance of research. The industrialmembers of Consortium are the companies-pro-ducers and consumers and investigation partners:five universities and such organizations as Army,Air Force, Navy, NIST, NASA. The developmentand implementation of this model was supportedby the state. The state Ohio allotted the many-million grant for realization of this project.

If the aim of consortia is to solve the strategicand organizational tasks on development of thenew technology, the centers of development andimplementation of industrial technologies are thebasic performers of this project. These organiza-tions should be recognized as the world classleaders in their field, equipped with the innova-tive equipment and having the highly qualifiedpersonnel. The example of such center in the fieldof materials joining is the EWI. It closely coop-erates in the work with research universities andindustrial sector which allows realizing of inno-vation developments and their successful imple-mentation into production. Since 1984 the Insti-tute has a constant state support according tothe Ohio Edison Program. The constant devel-opment, efficiency of work and high level of re-turn of investments attract private investors. In2010 the volume of private investments into thedevelopments of the EWI 20 times exceeded thevolume of state financing [7].

The innovation model of development and im-plementation of technologies proposed by theEWI was approved by the government of theUSA. The National Institute of Standards andTechnologies at the Trade Ministry accepted thenew state program on its basis to support thedevelopment of technological innovations in theUSA «Advanced Manufacturing TechnologyConsortia» in 2011. The budget of the programfor 2012 amounted to 12 Million USD. It envis-ages the support of development of such innova-tion directions as robotic technologies, nanoma-terials, new progressive materials, new progres-sive production technologies. In total, the stateallotted 75 Million USD to support the innova-tion programs in 2012 [8].

Table 4. Scheme of interaction of the purposeful industrial consortia and centers of development and implementation of industrialtechnologies

Purposeful industrial consortia

Centers of development and implementation of industrial technologies

Automation CastingAssembly ofelectronics

Stamping Control JoiningAdditive

technologiesTreatment

Production of metal for aircraft industry usingadditive technologies

× × × × ×

Decrease in automobile mass × × × × ×

High-speed assembly of batteries × × × ×

Ecologically clean production of electronics × × ×

Manufacturing of equipment for nuclear powerplants

× × × × ×

Automation of process of production of equip-ment for heavy-machine building

× × × × ×

1. Bucher, K. US competitiveness ranking continues tofall: Emerging markets are closing the gap. www.we-forum.org

2. (2011) Strengthening manufacturing competitiveness:Report from Conference 2010 on the Future of Mate-rials Joining in North America. EWI, Febr. 2.www.ewi.org

3. Conrardy, C. Materials joining and technology.www.weldingandgasestoday.org

4. State of the welding industry report: executive sum-mary. In: Weld-Ed. www.weld-ed.org

5. 2012 Global R&D Funding Forecast. In: Battelle.The Business of Innovation. www.battelle.org

6. Dezhina, I. (2011) Support of fundamental science inUSA: Lessons for Russia. Bytie Nauki, 94, 6—7.

7. Revitalizing American’s manufacturing innovation in-frastructure. Response to the NIST AMTech Requestfor Information. EWI. www.ewi.org

8. President Obama Launches Advanced ManufacturingPartnership. www.nist.go

52 11/2012

EXPERIENCE OF APPLICATIONOF HF ELECTRIC WELDING APPARATUS EK-300M1

IN SURGERY

O.N. IVANOVA1, A.T. ZELNICHENKO2, D.D. KUNKIN1, V.V. PEREKREST3 and V.A. TODORENKO3

1International Association «Welding», Kiev, Ukraine2E.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

3NTUU «Kiev Polytechnic Institute», Kiev, Ukraine

Presented are the data illustrating advantages and areas of application of HF electric welding apparatusEK-300M1 in surgery.

Keywo r d s : electric welding, soft live tissues, high-frequency apparatus EK-300M1, area of application, ad-vantages

Development of equipment and technology forhigh-frequency electric welding (HFEW) of softlive tissues became a priority for Ukrainian sci-entists, engineers and doctors.

In 1991 Prof. B.E. Paton and Prof. V.K. Le-bedev put forward a hypothesis and proposedmathematical model which were proved by seriesof experiments on animals. This allowed statinga possibility of electric welding of soft live tis-sues with preservation of their viability. It istime to develop the laboratory samples of tech-niques and experimental substantiation of thetechnology [1].

For this purpose an international project wasdeveloped in 1996 under the leadership ofProf. B.E. Paton with assistance of the E.O. Pa-ton Electric Welding Institute (PWI), A.A.Shalimov National Institute of Surgery andTransplantology, International Association«Welding» and commercial credit company«Consortium Service Management Group Tech-nologies Inc.», as well as series of medical insti-tutions of Ukraine. Common work was success-ful. The first variants of system with dosed powersupply, prototypes of samples of power sourcesand surgical instruments were proposed at thebeginning.

Development of the power sources was step-wise. The first experimental variant was devel-oped in 1992, the second in a period from 1995to 1996, and the third one in 2003. EK-300M1power source was developed in 2004 and its up-dated variant in 2007—2008 together with RI ofApplied Electronics of NTUU «Kiev PolytechnicInstitute» [2] and plant «Schyotmash» (Lubny,

Poltava region). Experimental work was kept onduring indicated periods in the PWI laboratoriesas well as in the clinics with assistance of thespecialists of engineer and medical orientation.

The first State Certificate about registrationof HFEW apparatus EK-300M1 in Ukraine wasissued in January, 2001. This is a date of begin-ning of practical application for new method oftissue saving HFEW technology in surgery. Thefurther State Registration Certificates were is-sued in 2004 and 2010. The State Certificatesabout registration in the Russian Federation(2006) and in Belarus (2009) were also obtained.Technology, method of HFEW and welding in-struments were patented in Ukraine, Russia,USA, European Union, Canada, China, Japanand Australia [3—15]. All these allowed proceed-ing to the application of method of soft live tissuewelding in different spheres at the clinics of 16regions of Ukraine as well as in Russia, namely,three Moscow and Saint-Petersburg clinics.

Medical instruments being a constituent ofthe HFEW complex is developed simultaneouslywith work over the power source. Such mainparameters as dimensions, form, weight of elec-trodes as well as requirement to design, i.e. com-fortable operation, access to place of HFEW,workability of instrument in manufacture andrepair are determined. All the types of electricwelding medical instruments (forceps, clampsand laparoscopes) are the instruments of bipolartype. Instruments used for endoscopic and thora-coscopic surgical procedures are of a great inter-est. Experience of manufacture and practical ap-plication of such type of the instruments is ac-cumulating. At that the experimental work isproceeded with and positive results are immedi-ately transferred in the clinical conditions.

© O.N. IVANOVA, A.T. ZELNICHENKO, D.D. KUNKIN, V.V. PEREKREST and V.A. TODORENKO, 2012

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Large scope of experimental and researchwork promoted accumulation of experience ofclinical application of HFEW and correspondingequipment.

The new electric welding technology is con-fidentially introduced in practice of surgerytreatment of the patients with different patholo-gies. Continuous accumulation of experience inthis direction allowed creating atlas «Tissue sav-ing HFEW technology in surgery» in 2009. Allthe participants of development of this technol-ogy, i.e. specialists of engineer and medical ori-entation, became the authors of this book.

The atlas reflects the whole way from the firststeps of researches up to the practical progressof clinicians who mastered around 130 types ofsurgical techniques as for 2009.

Development of given technology, improve-ment of HFEW equipment and new surgical tech-niques, increase of number of users of this equip-ment is observed today. Data on clinical appli-cation given by users of EK-300M1 equipmentto new (after 2009) surgical techniques can bean example of this technology development.These are Odessa Regional Oncology Dispensary,Kiev Center of Endocrine Surgery, Institute ofPediatrics, Obstetrics and Gynecology of theNAMS of Ukraine, Donetsk Regional Anti-Can-cer Center, V.P. Filatov Institute of Eye Diseasesand Tissue Therapy, Zhitomir Regional Chil-dren’s Hospital. These organizations successfullyapply new techniques of surgery treatment usingEK-300M1 for:

• septoplasty in children, in particular, innooks;

• laryngeal cancer;• treatment of pathologies of thyroid gland;• organ saving operations on removal of tumor

on ovaris in woman;• bladder cancer;• stomach cancer and breast cancer;• metastatic lesions of liver (hepatic resection,

marginal biopsy of liver, left and right hemi-hepatectomia, trisegmentectomia and bisegmen-tectomia);

• regmatogenous retinal detachment for block-ing of retina rupture;

• endoresection of intraocular neoplasms;• microsurgery of trabecular apparatus of eye

and iris.Original instruments for work in cavity of

vitreous body of eye and technique for obtainingadequate devitalization of tumorous focuses thatreduces a risk of uncontrolled bleeding in endore-section of uveal melanoma are developed up to

present time. HFEW technology is used at en-doscopic, mainly laparoscopic operations in pe-diatric surgery for treatment of different patholo-gies.

Application of EK-300M1 apparatus in opin-ion of medical staff allows achieving:

• significant reduction of blood loss at opera-tive surgical intervention;

• minimization of thermal and mechanical in-jury of tissue that result in preservation of livingcells and faster regeneration of tissues in placeof coagulation at retention of functional activityof organ, including the possibility of preservationof reproductive function;

• change of scheme of performance of opera-tive intervention with achieving of more simpli-fied access to injured organ (experience of neuro-surgical and urological operations);

• possibilities of performance of tissue savingoperative interventions;

• reduction of time of operative interventions,i.e. time of patient being under effect of narcoticdrugs;

• reduction of period of postoperative reha-bilitation;

• decrease of time of patient hospital stay;• elimination of application of foreign suture

material;• reduction of number of necessary medical

instruments;• improvement of work conditions for surgery

team, lighten the work of surgeon, in particular,at nooks;

• elimination of infiltration formation;• reduction of postoperative pain.Thus, we already have excellent instrument

for fight against illnesses and method of allevi-ating human’s physical suffering.

The present technology develops as a livingorganism. Control of power complex is improvedand new design solutions for apparatus and in-struments are developed. These efforts are di-rected at quality improvement and increase ofscope of surgical operations performed. The newelectric welding complex has been already testedunder clinical conditions. It will be added tosurgeons’ arsenal in the nearest time, i.e. thepossibilities of fight against the illnesses and ren-der the qualified help to the patients are ex-panded. It should be noted that principle of op-eration of EK-300M1 power sources in improve-ment remain the same in accordance with ac-quired patents.

1. Paton, B.E. (2004) Electric welding of soft tissues insurgery. The Paton Welding J., 9, 6—10.

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2. Zelnichenko, A.T., Mironichev, A.V., Perekrest,V.V. et al. (2008) Hardware-software complex forinvestigation of electric welding process of live softtissues. In: Welding and Related Technologies intothe Third Millennium: Abstr. Kiev: PWI.

3. Paton, B.E., Lebedev, V.K., Vorona, D.S. et al.Method of bonding of vessels and other animal andhuman hollow organs and device for its realisation.Pat. 39907 Ukraine. Prior. 25.03.98. Publ.16.07.2001.

4. Paton, B.E., Lebedev, V.K., Vorona, D.S. et al.Method of bonding of soft biological tissues and fix-ture for its realisation. Pat. 44805 Ukraine. Prior.25.03.98. Publ. 15.03.2002.

5. Paton, B.E., Lebedev, V.K., Vorona, D.S. et al.Bonding of soft biological tissues by passing highfrequency electric current therethrough. Pat.2002/0091385 A1 US. Publ. 11.06.2002.

6. Paton, B.E., Lebedev, V.K., Lebedev, A.V. et al.System and method for control of tissue welding.Pat. WO 03/070284 A2. Prior. 13.02.2003.

7. Paton, B.E., Lebedev, V.K., Vorona, D.S. et al.Bonding of soft biological tissues by passing highfrequency electric current therethrough. Pat.6,562,037 B2 US. Publ. 13.05.2003.

8. Paton, B.E., Lebedev, V.K., Lebedev, A.V. et al.System and method for control of tissue welding.Pat. 2003/0158551 A1 US. Publ. 21.08.2003.

9. Paton, B.E., Lebedev, V.K., Lebedev, A.V. et al.System and method for control of tissue welding.Pat. 6733498 US. Prior. 19.02.2002. Publ.21.08.2003.

10. Paton, B.E., Lebedev, V.K., Masalov, Yu.A. et al.System and method for control of tissue welding.Pat. 6,733,498 B2 US. Publ. 21.08.2003.

11. Paton, B.E., Lebedev, V.K., Lebedev, O.V. et al. In-strument for bonding of soft biological tissues of ani-mals and human. Pat. 72901 Ukraine. Prior.09.01.2004. Publ. 15.02.2006.

12. Paton, B.E., Lebedev, V.K., Vorona, D.S. et al. Bon-ding of soft biological tissues by passing high frequ-ency electric current therethrough. Pat.2005/0234447 A1 US. Publ. 20.10.2005.

13. Paton, B.E., Lebedev, V.K., Lebedev, O.V. et al.Method of welding of animal and human soft tissues.Pat. 75342 Ukraine. Prior. 19.06.2002. Publ.17.04.2006.

14. Paton, B.E., Lebedev, V.K., Lebedev, A.V. et al.Method of welding of biological tissue and devicefor welding of biological tissue (variants). Pat.77064 Ukraine. Prior. 13.02.2003. Publ. 15.10.2006.

15. Paton, B.E., Lebedev, V.K., Lebedev, A.V. et al.Method of welding of animal and human soft tis-sues. Pat. 2294171 RF. Prior. 19.06.02. Publ.27.02.07.

NEWS

REGULATOR FOR RESISTANCE WELDING

Regulator for resistance welding RKSM designed forcontrol of welding cycle in multispot AC resistancewelding machines was developed by «Obert» Ltd(Kiev) being specialized on development and manu-facture of electric automated mechanisms for weldingmachines.

Capability of setting of up to 99 welding modeswith their cyclic change in a process of operation ofwelding machine is the peculiarity of the regulator.

Immediately 8 valves and up to 32 valves usingexternal expander of the outlets can be controlled bythe welding cycle regulator.

Main technical parameters of regulatorPower consumed, V⋅A ........................... not more than 30Dimensions, m .................................... 0.16 × 0.16 × 0.29Weight, kg ........................................... not more than 4Number of welding current pulses ............................ 1—99Regulation of welding cycle time periods(compression—welding—pause between the pulses ofwelding—cooling—forging), s .................................. 0—399Regulation of real values of welding current, % ......... 0—99Duration of welding current rise (modulation)is programmed in the ranges, % ............................... 0—99

Regulation of load power factor .......................... 0.2—0.8Parameters of hyristor activation pulses:

voltage, V ........................................................... 24duration of pulse, μs ................................... 200 ± 150

Parameters of signal for power supply of outputDC devices:

voltage, V ...................................................... 24 ± 2current, A ...................................... not more than 0.4

Number of actuating devices, pcs ........................... up to 8Number of outputs (including pedal) ............................. 4

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SCIENTIFIC-TECHNICAL CONFERENCE«MODERN PROBLEMS OF METALLURGY,

TECHNOLOGY OF WELDING AND SURFACINGOF STEELS AND NON-FERROUS METALS»

On October 25—26, 2012, the Scientific-Techni-cal Conference «Modern Problems of Metal-lurgy, Technology of Welding and Surfacing ofSteels and Non-ferrous-Metals», devoted to the100th anniversary of Prof. I.I. Frumin and Prof.D.M. Rabkin, two famous scientists in the fieldof metallurgy and technology of welding and sur-facing of steels and non-ferrous metals, was heldat the E.O. Paton Electric Welding Institute ofthe NAS of Ukraine (PWI) in Kiev. The Con-ference was organized by PWI, Society of Weld-ers of Ukraine and International Association«Welding».

More than 100 scientists and specialists in thefield of welding and allied processes from re-search institutes, higher educational institutionsand enterprises of Ukraine, Russia and Polandparticipated in the work of the Conference.

The plenary session of the Conference wasopened by Prof. L.M. Lobanov, academician ofthe NAS of Ukraine. He described the course oflife of Prof. Frumin and Prof. Rabkin.

Isidor I. Frumin was given an employment atthe PWI by Evgeny O. Paton in 1937. Since 1941to 1945 he took part in the Great Patriotic War.He finished the war in Berlin having a militaryrank of major and being the officer of chemicalservice of anti-aircraft division.

After demobilization, I.I. Frumin returned tothe PWI, where he headed the chemical and fluxlaboratories, and then he was a chief for almost30 years of the Department of physical-metallur-gical problems of surfacing the wear-resistant andheat-resistant steels. Together with his team heaccomplished a complex of fundamental worksin the field of metallurgy of welding, theory offormation of pores and cracks in welding, devel-oped and implemented the first fused fluxes inthe industrial production.

A great contribution was made by I.I. Fruminto the creation of scientific and practical basesof the mechanized surfacing. Under his supervi-sion the first flux-cored wires for surfacing, newmethods and technologies of surfacing were de-veloped which found the wide application in dif-ferent branches of industry. I.I. Frumin became

the first laureate of Evgeny Paton Prize of NASUfor research, development and implementationof the mechanized surfacing of hot milling rolls.The State Prize of the USSR was awarded to himin the team of scientists for the development offlux-cored wires for welding and surfacing.

Daniil M. Rabkin began to work at the PWIin 1939 after graduation from the Kiev IndustrialInstitute (now NTUU «Kiev Polytechnic Insti-tute»). In 1941—1943 he joined the ranks of RedArmy. In 1943 he was recalled from the front tothe PWI for urgent solution of problems con-nected with the development and implementa-tion of technologies of welding of armor struc-tures and shells at the plants of Ural and Siberia.All his further activity is connected with thePWI where he was dealing with the problems ofmetallurgy of welding of light alloys.

D.M. Rabkin showed himself as a talentedresearcher of processes of melting of aluminiumalloys in arc welding and physical-chemical re-actions in arc and weld pool. He fulfilled thefundamental studies both in the field of metal-lurgy and materials science of aluminium alloysand also in the development of new technologiesof their welding, such as mechanized semi-openarc using halogenide fluxes, electroslag, electronbeam, etc. He was awarded the Evgeny PatonPrize of NASU for the monograph «Metallurgyof fusion welding of aluminium and its alloys».

Then, Prof. I.A. Ryabtsev, the Chief of PWIDepartment of physical-metallurgical problemsof surfacing the wear- and heat-resistant steels,made a speech. He said, in particular, that thoseinvestigations and traditions are continued andfurther developed in the Department which wereestablished by I.I. Frumin. Over the recent yearsthe Department is dealing with research and de-velopment of the new methods of optimizationof structure and properties of the deposited met-al. For this, the effect of structural heredity isused. To realize this effect, the nano-sized carbidecompositions, which have an influence on struc-ture and properties of the deposited metal with-out change of its chemical composition, are addedinto composition of surfacing flux-cored wires.

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In collaboration with the S.P. Timoshenko In-stitute of Mechanics of NASU the Departmentis dealing with the development of mathematicalmodels and methods of calculation of structuraland stress-strain states of plane and cylindricalparts in surfacing and service under the condi-tions of simultaneous action of cyclic thermaland mechanical loads. Mathematical modelsmake it possible to evaluate the service life ofdeposited mill rolls, dies and other similar partsusing the method of calculation. In collaborationwith the PWI Department of mathematical in-vestigations an expert system «Naplavka» (Sur-facing) was developed. The development of thissystem allowed systematizing the comprehensiveknowledge on surfacing materials, technologiesand technique of surfacing of almost all the partswhich are subjected to surfacing in CIS countries.Taking into account the present feasibilities inobtaining the practical knowledge and experi-ence on surfacing, the expert system can be usedsuccessfully in the process of education at thehigher educational institutions.

Prof. A.Ya. Ishchenko (PWI), Corr.-Memberof NASU, delivered a paper «Challenging tech-nologies of welding the high-strength aluminiumalloys». Aluminium alloys of different systemsof alloying are widely used in aircraft, aerospaceand defense engineering. In the paper the physi-cal-metallurgical processes, proceeding in theirwelding, such as formation and measures of pre-vention of inclusions of oxide film in weld metal,causes and measures of prevention of porosity inweld metal in fusion welding, peculiarities ofcrystallization of multi-component alloys, forma-tion and measures of prevention of hot cracks,chemical and structural heterogeneities, wereanalyzed. Characteristic of new and updatedmethods and technologies of welding using anelectric arc, electron beam and laser source ofheating was given. Over the recent years, thecomplexly-alloyed aluminium alloys with mi-croadditions of scandium and zirconium havebeen developed, which are characterized by thehigh manufacturability and strength. Their weld-ability by fusion using updated technologies ofwelding is characterized as good or satisfactory,and the tensile strength of heat-hardened de-formed semi-products reaches 750 MPa. The im-provement of characteristics of weldability ofthese materials will provide the progressive de-sign of many new products of transport purpose,such as airbuses, cars of high-speed trains, prod-ucts of defense purpose, that will increase thetechnical and economic characteristics of theirproduction and service.

Paper of Dr. E.F. Pereplyotchikov (PWI) wasdevoted to the achievements of the Institute inthe field of plasma-powder surfacing (PPS), theprogress of which is indispensably connectedwith the name of Frumin I.I. Under his supervi-sion, the integrated and purposeful investigationsof technological features of plasma surfacing, de-velopment of surfacing powders and equipment,as well as implementation of the process in thedifferent branches of industry were carried out.PPS is especially effective under the conditionsof manufacture of different-purpose reinforce-ment with sealing surfaces, deposited by alloyson cobalt, nickel, iron, copper base. A large ex-perience is gained in surfacing of parts as smallgates, stop valves (DN50), and also large ones(DN1000 and larger) for stationary and transportpower plants, chemical enterprises, oil-and-gaspipeline, Technology and surfacing powder PR-X18FNM for PPS of worms of extruders of poly-meric machines, operating under the conditionsof abrasive wear and corrosion effect of environ-ment, developed at the PWI, represent a greatinterest. The plasma surfacing in repair of wormsis especially effective, as it allows at small ex-penses not only restoration of expensive part, butalso increase its serviceability, and, moreover,restoration of worn-out worms several times. PPSis used in industry for surfacing of exhaust valveof internal combustion engines of different typesand sizes, staring from valves of motor cars of20—35 mm diameter and finishing by valves ofship diesels of 300—450 mm diameter. In surfacingof valves the most important advantage of PPSis greatly manifested, i.e. possibility of deposi-tion of thin layers at low thermal effect on theparent metal.

The paper of Dr. O.G. Kuzmenko (PWI) de-scribed the developments of the Institute in thefield of electroslag surfacing (ESS). A uniquedesign of a non-consumable electrode, i.e. cur-rent-carrying mould (CCM), was made at theInstitute. With use of the CCM the filler mate-rials can be fed to a slag pool in the form of atubes, bars, wires, shot, liquid filler material,etc. The grainy material is most promising forsurfacing in CCM. With its use it is possible toproduce the deposited layers of preset sizes andchemical composition, as well as to influence ac-tively on the processes of deposited metal crys-tallization and its properties. The largest expe-rience was gained in manufacturing and restora-tion surfacing of mill cast iron rolls using a shot.At PWI a method of producing the multi-layermetal by ESS with liquid metal was developedand implemented in industry. The process is

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started with pre-fusion of a hard billet, then themetal, melted in a separately-standing furnace,is poured on the pre-fusion surface and the elec-troslag process is maintained by non-consumablegraphitized electrodes. The deposited metal issolidified gradually by decreasing the electroslagprocess power. Owing to this, the metal is wellrefined and solidified in upward direction with-out defects of a shrinkage nature. At the presenttime, the developed technology passed successfultest in industry for producing billets of bimetaldies, restoration of dies of hot deforming, etc.

Engineer A.Yu. Pasechnik (Donetsk NationalTechnical University) described the works car-ried out at the laboratory of welding and surfac-ing in DNTU. The laboratory is dealing with thedevelopment and implementation of technologiesof repair, strengthening and manufacturing ofparts and components of mining and metallurgi-cal equipment using the electroslag process. In-dispensable specifics of the developed technolo-gies is the possibility of their industrial realiza-tion directly in the sites of the equipment service,and also the use as initial materials of metalwastes.

Prof. V.Yu. Konkevich (All-Russian Instituteof Light Alloys, Moscow) presented the paperabout technologies of production and applicationof granulated aluminium alloys. The main advan-tages of a granular production of aluminium al-loys consists in a feasibility of using the simplertechnological scheme of producing the thin-walled semi-products; providing the economicalproduction of products due to reduced cycle incombination with a high yield of an efficientmetal; producing of semi-products of complexly-alloyed alloys, containing components in theircomposition in the amounts exceeding their lim-iting solubility in an equilibrium state.

Prof. E. Turyk (Institute of Welding, Gli-wice, Poland) presented the paper «Experimen-tal investigations of thermal resistance and resis-tance at cyclic temperature and constant me-chanical loads of metal used for surfacing of roll-ers of machines for the continuous casting of bil-lets». These experimental investigations werecarried out using the metal deposited by wireswhich are used at Polish metallurgical plants forrestoration and strengthening of rollers ofMCCB. For comparison, the base metal of rollers(steel 34KhM) was tested using the same proce-dures. As a result of investigation, it was foundusing both procedures that during these testsspecimens of steel 34KhM had the best proper-ties. Among the deposited specimens the austeni-tic deposited metal Kh18N10 had the best prop-

erties. Properties of metal of martensitic 10Kh13and martensitic-austenitic Kh13N4 classes hadthe lower properties. The lower resistance ofspecimens of deposited metal of steel typesKh18N10, 10Kh13 and Kh13N4 is, probably, con-nected with precipitations of chromium carbidesat the boundaries of grains.

The paper of Dr. V.N. Matvienko (PriazovskyState Technical University, Mariupol, Ukraine)described the problems of increasing the life ofmill rolls using surfacing and arc metallization.The work was fulfilled in collaboration with theIlljich Metallurgical Works. Because of increas-ing the cost of mill rolls, as well as due to growthin prices for surfacing materials, not only resto-ration of rolls, but also the development andmastering the production of surfacing materialsdirectly at the Works become of current impor-tance. At present, the production of alloyed sur-facing strip 20Kh4MFB has been mastered at theWorks for surfacing of mill rolls. The applicationof this strip (alongside with steel strips 08(rimmed), 20 (semi-killed)) in combination withfused or agglomerated fluxes make it possible tosurface the layers, the mechanical and serviceproperties of which correspond to the service con-ditions of rolls. The surfacing of mill rolls usingthe surfacing materials manufactured at theIlljich Metallurgical Works provides a lowerfraction of expenses (33—45 %) for restoration ofworn-out rolls as compared with the cost of thenew ones.

The paper of Dr. K. Madej (Institute of Weld-ing, Gliwice, Poland) included the problems ofwelding of as-heat-treated structural steels ofhigh strength with 690—1100 MPa yield strength.The brief characteristic of these steels, purpose,chemical composition, mechanical propertieswere given. Peculiarities of technology of weld-ing of high-strength steels, in particular, the ef-fect of heat input of arc welding on structure andmechanical properties of welded joints were con-sidered. One of the main defects in welding ofsteels of similar type is the cracks, and the paperpresented data on causes of their initiation andpossible measures of their prevention.

Dr. A.G. Poklyatsky (PWI) presented paperat the Conference about the efficiency of appli-cation of friction stir welding for producing thepermanent joints of aluminium alloys. The for-mation of welds in solid phase prevents the for-mation of pores, macroinclusions of oxide film,hot cracks and other defects. The absence of arcdischarge and molten metal allows producing thepermanent joints without application of shield-ing gas, as also avoiding the ultra-violet radiation

58 11/2012

of arc, fumes, burn-out of alloying elements. Thereduction of heat effect on metal promotes thedecrease in residual stresses and deformations,thus leading to the lower distortion of weldedstructures and increasing their resistance to frac-ture. FS-welded joints of aluminium alloys havea large resistance to initiation and propagationof service cracks and a high fatigue strength.

In total, 27 papers and presentations weremade at the Conference plenary session. In ad-dition, more than 10 poster papers were presentedin the reading-hall of the PWI library, repre-senting a great interest to the Conference par-ticipants.

Prof. Ryabtsev I.A., PWI

NEWS

SPHERICAL TUNGSTEN CARBIDE

Technology for manufacture of powder refractory ma-terials, in particular cast tungsten carbides WC +W2C with spherical-shaped granules, using thermalcentrifugal spraying was developed at the E.O. PatonElectric Welding Institute. Special unit SFERA-2500was developed for centrifugal thermal spraying allow-ing obtaining of the granules from 50 up to 850 μmsize at 15—20 kg material per hour efficiency.

Granulated tungsten carbides with HV0.1 ≥≥ 3000 MPa hardness exceeded the similar materialson its physical-mechanical and technological proper-ties, and are successfully used for increase of wearresistance of the parts of drilling equipment and tool.

Developed material is widely used for plasma-pow-der, laser and oxy-acetylene cladding as well asstrengthening of parts by impregnation method.

Strip cladding material of 8.0 × 3.0 mm sectionwas developed on the basis of fused tungsten carbide

powders. It is delivered in the bundles for automaticplasma cladding or rods for oxy-acetylene and atomic-hydrogen cladding of parts of drilling equipment.

Unit for thermal centrifugal spraying SFERA-2500

Drill pipe lock cladded by spherical tungsten carbide

General view of spherical granules of fused tungsten carbides

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