CONTENTS

SCIENTIFIC AND TECHNICAL

Zyakhor I.V. and Kuchuk-Yatsenko S.I. Friction weldingof PIM heat-resistant steel to steel 40Kh ................................. 2

Grinberg B.A., Elkina O.A., Patselov A.M., InozemtsevA.V., Plotnikov A.V., Volkova A.Yu., Ivanov M.A., RybinV.V. and Besshaposhnikov Yu.B. Problems of stirringand melting in explosion welding (aluminium—tantalum) ............ 12

Khorunov V.F., Maksymova S.V. and Stefaniv B.V.Effect of palladium on structure and technologicalproperties of Ag—Cu—Zn—Ni—Mn system brazing filleralloys ..................................................................................... 20

Knysh V.V., Barvinko A.Yu., Barvinko Yu.P. and YashnikA.N. Substantiation of «leak-before-break» criterion forvertical cylindrical tanks for oil storage .................................... 26

Lebedev A.V. Transistor power sources for electric arcwelding (Review) .................................................................... 30

INDUSTRIAL

Tsaryuk A.K., Ivanenko V.D., Skulsky V.Yu., MoravetskyS.I., Gavrik A.R., Strizhius G.N., Nimko M.A., MazurS.I., Trojnyak A.A., Odin Yu.V., Derkach O.V. and KuraR.I. Technology of repair welding of boiler unitassemblies without postweld heat treatment ............................ 37

Senchishin V.S. and Pulka Ch.V. Modern methods ofsurfacing the tools of agricultural tillers and harvesters(Review) ................................................................................. 43

Goloborodko Zh.G. Experience in hardfacing ofpropeller shafts at the PJSC Kherson Shipyard ........................ 50

Kuzmenko O.G. Effect of flux composition onthermal-physical and physical-chemical processes inliquid-metal electroslag surfacing ............................................ 52

Developed at PWI ................................................................... 25

© 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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Yu.N.Lankin I.K.PokhodnyaV.N.Lipodaev V.D.Poznyakov

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September2012

# 9

FRICTION WELDING OF PIM HEAT-RESISTANT STEELTO STEEL 40Kh

I.V. ZYAKHOR and S.I. KUCHUK-YATSENKOE.O. Paton Electric Welding Institute, NASU, Kiev Ukraine

Experimental data are given on evaluation of structure of heat-resistant steel AISI310 produced by thepowder injection moulding technology. The investigation results are presented on peculiarities of formationof dissimilar joints between steel AISI310 and structural steel 40Kh under different thermal-deformationcycles of friction welding in manufacture of bimetal shafts for automotive engine turbocharger rotors.

Keywo r d s : friction welding, bimetal joints, pow-der injection moulding, welded joints, turbocharger ro-tor shafts

One of the new powder metallurgy methods ispowder injection moulding [1—5], which in theEnglish technical literature is collectively knownas the PIM-technology. Powder injection mould-ing has been gaining an increasingly wider ac-ceptance in the last years owing to a number ofadvantages over traditional methods of metalprocessing, first of all in manufacture of com-plex-geometry and mass-production parts. Ac-cording to the data given in [3], density of thefinished parts produced by the PIM-technologyis 96 to 100 % of its theoretical value, while thedetected pores and non-metallic inclusions havesmall sizes and a spherical shape, and are uni-formly distributed in the bulk.

The promising market for the parts producedby the PIM-technology is the automotive engineconstruction industry. The issue of current im-portance in terms of technology and economy isapplication of the PIM-technology to manufac-ture complex-configuration parts. For example,these are wheels of bimetal shafts for automotiveengine turbocharger rotors (TCR). Compared tothe currently applied investment casting method,the PIM-technology provides an increased pro-ductivity, minimal possible deviations of sizesand a high quality of the surface of the wheelsfor the TCR shafts.

The technological cycle of manufacture of thebimetal TCR shafts provides for the use of fric-tion welding (FW) of a heat-resistant alloy wheelto a structural steel shank. Friction welding isapplied to advantage to join materials producedby the casting, thermo-mechanical deformationand powder metallurgy methods [6—8]. However,no information on application of FW for the partsproduced by the PIM-technology has been foundin technical literature so far. Therefore, it is ofscientific and practical interest to study the effectof structure of the PIM-materials on the possi-bility of joining them to structural steel to manu-facture bimetal TCR shafts.

The purpose of this study was to investigateformation of dissimilar joints between heat-resis-tant steel AISI310 produced by the PIM-tech-nology and structural steel 40Kh under differentthermal-deformation cycles of FW for manufac-ture of bimetal TCR shafts.

General view of the TCR shaft wheels madeby the PIM-technology from steel AISI310 (feedstock – BASF «Catamold» [3]) is shown inFigure 1. Chemical composition of the materialswelded, as well as their mechanical propertiesare given in the Table.

Austenitic stainless steel AISI310 (domesti-cally produced analogue – steel 20Kh25N20S2(EI283)) combines satisfactory heat resistanceand high oxidation resistance at high tempera-tures. Bimetal TCR shafts were produced by join-

Chemical composition and mechanical properties of materials welded

Steel gradeContent of elements, wt.% Mechanical properties

C Cr Nb Si Mn Fe Ni σy, MPa σt, MPa δ, % ψ, %

AISI310 515 >40 >50

40Kh 0.36—0.40 0.8—1.1 — 860

ing wheels of steel AISI310 to shanks of struc-tural steel 40Kh under different conditions ofconventional and combined FW [7, 8].

Experiments on FW were carried out by usingmachine ST120, which was upgraded to imple-ment different thermal-deformation cycles corre-sponding to conventional, inertia and combinedFW [9]. Structure of the materials welded andof the bimetal joints was examined by opticalmicroscopy («Neophot-32», Germany) and byscanning electron microscopy (SEM) (JSM-35CA, JEOL, Japan). X-ray spectrum microana-lysis (EDS-analyser INCA-459, «Oxford Instru-ments», Great Britain; probe diameter – ap-proximately 1 μm) was conducted, and micro-hardness of the joining zone metal was measuredunder a load of 1—5 N (microhardness meterM400, LECO, USA).

Parameters of the FW process were variedwithin the following ranges: heating pressurePh = 50—150 MPa, forge pressure Pf = 150—300 MPa, peripheral velocity v = 0.5—2.5 m/s,

heating time th = 5—30 s, deceleration time td == 0.2—2.5 s, and burn-off rate vb = 0.1—1.0 m/s.Diameter of the billets welded was 15 mm.

Microstructures of the base metal of steelAISI310 and compositions of structural compo-nents are shown in Figures 2 and 3. Steel AISI310has equiaxed crystalline structure with grain sizeof about 60 μm. Optical microscopy allows re-vealing dark and light grains (Figure 2, a), in-dividual pores with a size of up to 15 μm locatedalong the grain boundaries, and randomly ar-ranged particles of non-metallic SiO2 inclusionswith a size of 2—10 μm (Figure 3, spectrum 4).

Structure of the dark grains (see Figure 3,spectra 1 and 8) is lamellar, consisting of alter-nate light- and dark-etchable laminae less than1 μm thick. The content of chromium in the grains

Figure 1. Wheels for TCR shafts of steel AISI310 producedby the PIM technology (a), and welded TCR shaft (b)

Figure 3. Microstructure of the base metal of steel AISI310(SEM), and results of X-ray spectrum microanalysis of me-tal of the investigated regions (wt.%)

Figure 2. Microstructures of steel AISI310 (optical microscopy)

9/2012 3

with a lamellar structure is somewhat lower,compared to its content in the light grains.Chemical composition of the light grains (spec-trum 2) corresponds to the requirements of stand-ard AISI, except for the increased niobium con-tent. Dispersed (1—2 μm) particles of a roundshape (spectra 5 and 6) can be distinguished inthe bulk of the light grains. These particles donot differ in chemical composition from those inthe bulk of the dark grains. However, they havelower niobium content. No substantial liquationof alloying elements and impurities along thegrain boundaries is fixed (spectra 3 and 7). Pres-ence of the grains with a lamellar structure evi-dences that the sintering process was performedat a temperature close to TS [10—12].

Formation of the dissimilar joints betweensteels AISI310 and 40Kh was investigated underdifferent thermal-deformation cycles correspond-ing to conventional and combined FW.

Mode 1 («soft» mode) was conventional fric-tion welding (CFW) [7, 8] with the minimalvalues of Ph, Pf and td, and the maximal valuesof v and th in the ranges under investigation.Mode 2 («rigid» mode) was CFW with the maxi-mal values of Ph and Pf, and the minimal values

of v, th and td in the ranges under investigation.Mode 3 was the technology developed by theE.O. Paton Electric Welding Institute for com-bined FW with controlled deformation. For thistechnology the values of v and Ph were set pro-ceeding from the requirement to provide the cer-tain deformation rate (burn-off rate) in heating,which was varied during the heating processwithin the vb = 0.1—1.0 m/s range, forge pressurebeing applied at a stage of rotation decelerationregulated following the preset program [9]. Incombined FW the values of th and td were seton the basis of the results of preliminary experi-ments, so that the total length loss for all themodes investigated was Δw = 6 mm.

As seen from Figure 4, a, deformation of thebillets during welding in mode 1 occurs mainlydue to steel 40Kh. The transition zone with awidth varying from 25 μm in the peripheral partof a section to about 500 μm at the centre, which

Figure 4. Macrosection of the AISI310—40Kh steel joint made in mode 1 (a), panoramic view of section of the weldedjoint (b), and welded TCR shaft after tensile tests (c)

Figure 5. Microstructure of the AISI310—40Kh steel joint(mode 1)

Figure 6. Distribution of chromium (a) and nickel (b) acrossthe joining zone between steels AISI310 and 40Kh (mode 1)

4 9/2012

is non-uniform across the sections of the billets,can be seen within the joining zone (see Fi-gure 4, b).

X-ray spectrum microanalysis fixed a variablecomposition of the transition zone across thewidth (see Figures 5 and 6). The presence ofuniformly distributed non-metallic SiO2 inclu-sions in the transition zone proves that the latterforms on the side of steel AISI310. Size of theSiO2 particles in the transition zone is 3—4 μm,this being indicative of their partial dissolution.Monotonous growth of the concentration of chro-mium and nickel in transition from steel 40Khto steel AISI310 (Figure 6) may be related todiffusion migration of these elements during acomparatively long stage of friction heating (th == 30 s). Interlayers of an intermediate composi-tion (Figures 5 and 6), having an increased hard-ness (Figure 7), can be seen in the transitionzone on the side of steel 40Kh. Formation of thiscomposition of interlayers may result from dilu-tion of contact volumes of metal of the steelswelded in the liquid or solid-liquid state at theinitial stages of FW. These interlayers lead todeterioration of ductility and decrease in corro-sion resistance of the joints [13, 14].

Deformed, radially elongated grains can beseen in the zone of the thermal-deformation effecton the side of steel AISI310 (see Figure 7), re-sidual porosity of the base metal being conservedin this zone. Steel AISI310 has a fibrous structurewith structural components up to 10 μm in sizein the immediate proximity to the transition zone.No segregations of non-metallic SiO2 inclusionsand porosity were detected in this zone.

Annealing and repeated etching of the weldedjoints were carried out to reveal structure of thetransition zone. As a result, the transition zonewas found to have a lamellar structure and acomposition intermediate between the steelswelded (Figure 8, spectra 3—8).

In tensile tests, fracture of the welded jointsoccurred in the joining zone (see Figure 6, b).Chemical composition of metal on both sides ofthe fracture was close to that of steel AISI310.Fractography of the fractures revealed the pres-ence of the two types of fracture within the limitsof a section, i.e. tough and quasi-brittle (Fi-gure 9). The increased niobium content was de-tected over the entire fracture surface and, par-ticularly, in the circumferential region of thequasi-brittle fracture (see Figure 9, spectrum 2).

Figure 7. Microstructures of the transition zone (a) (indentations HV3, MPa⋅10—1) and metal of the joint on the sideof steel AISI310

Figure 8. Microstructure of the AISI310—40Kh steel jointmade in mode 1 after annealing, and results of X-ray spec-trum microanalysis of metal of the investigated regions(wt.%)

9/2012 5

Results of the fractographic examinationsshow that fracture of the joints occurs betweenthe transition zone and steel AISI310, and is lo-calised in segregation clusters of the redundantphases with the increased niobium content.

Therefore, the presence of an insignificant po-rosity and dispersed inclusions of SiO2, havinga high melting point (Tmelt = 1713 °C), in thebase metal of steel AISI310 does not exert anysubstantial effect on formation of the weldedjoints. At the same time, the presence of thesegregations of niobium, which forms eutectic(Tmelt = 1355 °C) with iron and leads to the«contact melting» phenomenon [15], in the basemetal exerts a negative effect on the compositionand mechanical properties of the joints in the«soft» mode of CFW.

Structure of the joint made in the «rigid»mode of CFW (mode 2, after annealing) is shownin Figure 10. The 40—60 μm wide transition zone,which is almost uniform across the section of thebillets and consists of alternate layers with dif-ferent etchability, was fixed in the joining zone.The character of variations in the concentrationsof chromium and nickel (Figure 11) when passingfrom steel 40Kh to steel AISI310 cannot be causedby diffusion migration of these elements, but isa result of dilution of the steels welded.

Analysis of microstructure of the weldedjoints made at th = 0.5—1.5 s (initial stage of theFW process) showed that the lamellar structureof the joining zone was formed at the early stagesof the FW process, and was determined by thecharacter of contact interaction of the faying sur-faces at the set level of the process parameters.At a low peripheral velocity and a high heatingpressure the predominant mechanism of contactinteraction at the initial stages of FW is the proc-ess of deep tearing out and dilution of contact

Figure 9. Fracture surface of the AISI310—40Kh steel joint(mode 1) on the side of steel 40Kh (a), regions of tough(b) and quasi-brittle (c) fractures, and results of X-rayspectrum microanalysis of metal of the investigated regions(wt.%)

Figure 10. Macrosection (a) and microstructure (b) of theAISI310—40Kh steel joint (mode 2)

6 9/2012

volumes of the materials welded in the plasticisedor solid-liquid state, going to a depth of severalhundreds of microns [16, 17].

Therefore, the key feature of the joints madein the «rigid» mode of CFW is the presence ofalternate constant-composition interlayers (Fi-gure 11), including those corresponding to steelsof the martensitic (see Figure 11, spectra 3 and7) and austenitic grades (Figure 11, spectra 2,4—6 and 8). It is commonly supposed [7, 8] thatFW is a solid-state process of joining of materials.However, the presence of the lamellar structureconsisting of alternate «alloys» of different com-positions allows a conclusion that the local melt-ing processes occurring in the contact interactionzone, at least at the initial stages of the FWprocess, play an important role in formation ofthe dissimilar joints.

No interlayers with the increased niobium con-tent were fixed in the joint made in mode 2,despite the presence of local clusters of this ele-ment in the immediate proximity to the joiningzone (see Figure 11, spectrum 9). The main dif-ference between the «rigid» and «soft» modes ofCFW consists in the burn-off rate during thefriction heating process (vb = 0.9 mm/s for mode2, and 0.15 mm/s for mode 1). It is likely thata high burn-off rate and a short time of the heat-ing stage (th = 6 s) in the «rigid» mode of CFWprevent formation of interlayers with the in-creased niobium content.

The joint made in mode 3 (combined FW withcontrolled deformation) was found to contain noincreased-hardness interlayers. The up to 40 μmwide transition zone, being almost uniform acrossthe section (Figure 12) and having a grain sizeof 5—6 μm, was fixed in this joint. The compo-sition of metal of the transition zone across itswidth corresponded to that of steel of the austeni-tic grade (Figure 13, spectra 2, 3, 5—7), thispreventing any decrease in corrosion resistanceof the joints and eliminating the risk of cracking.

Of notice is the absence of local segregationsof phases with the increased niobium content andthe dramatic change in the concentrations ofchromium and nickel in passing from steel 40Khto steel AISI310, this being indicative of mini-misation of dilution of the steels in the solid-liq-

Figure 11. Microstructure of the AISI310—40Kh steel joint(mode 2), results of SEM and X-ray spectrum microanalysisof metal of the investigated regions (wt.%)

Figure 12. Macrosection (a) and microstructure (b) of theAISI310—40Kh steel joint, mode 3 (indentations HV3,MPa⋅10—1)

9/2012 7

uid state and insignificant development of thediffusion processes in the joining zone. The ther-momechanically affected zone with fine-grained,dynamically recrystallised structure is immedi-ately adjacent to the transition zone. No poresand non-metallic inclusions of SiO2 were fixedin this zone, which was clearly shown by analysisof microstructure of the joint both in the as-welded state (see Figure 13) and after annealing(Figure 14). Disappearance of the SiO2 particlesfrom the joining zone may be related to theirpartial dissolution and the earlier establishedphenomenon of destruction of oxides by a flowof moving dislocations [18—21], including underthe thermal-deformation conditions of inertia andcombined FW.

Fracture of the joints in mechanical tests oc-curs in the base metal of steel AISI310 (Fi-gure 15). The joints had tensile strength at alevel of σt = 580—630 MPa. As shown by meas-urements of microhardness, metal of the joiningzone with the fine-grained, dynamically recrys-tallised structure is characterised by the increasedvalues of strength (see Figure 12), which is at-tributable to a substantial decrease (from 60 to5—6 μm) in size of structural components.

Comparative analysis of structure and chemi-cal composition of the zone of joints betweenPIM-steel AISI310 and steel 40Kh, made underdifferent thermal-deformation cycles of FW sug-gests the following mechanism of formation ofthe transition zone in FW of the investigatedmaterials combination.

At the initial stage of the FW process, becauseof a lower value of thermal conductivity of theaustenitic steel, the surface of maximal shear de-formations («friction plane») shifts towardsaustenitic steel AISI310. The «friction surfaceshift» phenomenon is known for other combina-tions of dissimilar materials as well [22—24]. Themaximal values of the heating temperature are

Figure 13. Microstructure of the AISI310—40Kh steel joint(mode 3), results of SEM and X-ray spectrum microanalysisof metal of the investigated regions (wt.%)

Figure 14. Microstructure of the AISI310—40Kh steel jointafter annealing (mode 3)

Figure 15. Welded TCR shaft after tensile tests (a), fracture surface in base metal of steel AISI310 made by thePIM-technology (b)

8 9/2012

fixed in the friction plane. The transition zone,which in fact is an AISI310 steel based alloy«deposited» on steel 40Kh, forms between thefriction plane and steel 40Kh. Burn-off of thebillets and, therefore, displacement of the tran-sition layer outside the section do not occur atthe initial stage of the heating process.

As duration of the heating stage increases,width of the transition zone in peripheral partsof the section remains almost unchanged, whereasin the central part of the section it grows due tofurther displacement of the friction plane to-wards steel AISI310. This is proved by metal-lographic examinations of the joints made at dif-ferent durations of the heating stage. As a result,the shape of the transition zone in the centralpart of the section becomes convex, directed tosteel AISI310.

When CFW is performed in the «soft» mode,in the friction plane, where the heating tempera-ture and tangential deformations are maximal,the rate of diffusion of alloying and impurityelements may be commensurable with that of themelts. As a result, clusters of low-melting pointphases, in particular of the eutectic phases of ironwith niobium, the redundant content of whichwas fixed in the base metal of steel AISI310,form in this zone.

The formed lamellar structure of the transitionzone persists in passing to the quasi-stationarystage of heating, which is accompanied by burn-off of the billets. The process of burn-off of thebillets occurs mainly due to steel 40Kh, the tran-sition zone metal being partially displaced out-side the section in the form of a thin layer de-posited on the reinforcement surface of steel40Kh. As the plane of maximal shear deforma-tions is located in the austenitic steel, restorationof the «deposited» layer takes place simultane-ously, i.e. there comes the state of equilibriumbetween the processes of formation and displace-ment of the transition zone metal.

The width and shape of the transition zoneare determined by a difference in thermal-physi-cal characteristics of the steels welded, and bythe friction process parameters. «Soft» mode 1is characterised by a low value of burn-off rate(about 0.15 mm/s). In this case the width ofthe transition zone is maximal, and the conditionsfor liquation of impurity elements and formationof low-melting point phases in the friction planeare most favourable. The applied force of forgingperformed on the non-rotating billets causes de-crease in thickness of the transition zone, as wellas partial displacement of the variable-composi-

tion interlayers and clusters of low-melting pointphases from the joint, the remainder determiningmechanical and service (corrosion, fatigue) prop-erties of the welded joint.

As the peripheral rotation velocity is de-creased and the heating pressure is increased(«rigid» mode 2 of CFW), the deeper and deeperlayers of metal of the billets welded are involvedin the shear (tangential and radial) deformationprocess, burn-off rate substantially grows (up to0.9 mm/s), and welding time is reduced. As aresult of such a «rigid» thermal-deformation cy-cle of FW, the narrow (up to 60 μm) transitionzone forms in the joint, consisting of alternatevariable-composition interlayers, including thosecorresponding to steel of the martensitic grade.These interlayers forming at the initial stages ofthe FW process are not fully displaced from thecontact zone at the subsequent stages of heating,and persist in the welded joint after forging.

As forging in CFW is performed after stoppingof rotation of the billets, the effect on the tran-sition zone metal is characterised by the presenceof the radial deformation component. Burn-offoccurs mainly due to deformation of metal in theheat-affected zone of steel 40Kh. Only decreasein thickness of the variable-composition interlay-ers takes place in this case.

The joints made by combined FW with con-trolled deformation (mode 3) are free from theabove imperfections of structure. The fact thatthe transition zone contains no alternate inter-layers corresponding in composition to marten-sitic and austenitic steels is caused by a peculiarquality of the initial stage of the FW processwith controlled deformation. The process pa-rameters (pressure, peripheral velocity) at thisstage were set based on the requirement to min-imise the processes of deep tearing out and dilu-tion of contact volumes of the materials welded.

The absence of local segregations of low-melt-ing point phases in the joint is caused by mini-misation of duration of the quasi-stationary stageof heating, which is achieved by providing thepreset rate of deformation (burn-off rate) of thebillets at a certain combination of pressure andperipheral velocity values.

The forging stage plays the key role in forma-tion of the joints. Programming of the durationof the rotation deceleration stage and applicationof increased forge pressure at this stage allow theburn-off rate and the intensity of deformation ofcontact volumes of the steels welded to be con-siderably increased. The fixed dramatic changein the concentrations of chromium and nickel

9/2012 9

within the joining zone in mode 3 is indicativeof minimisation of the processes of dilution ofthe steels welded and local melting in the contactzone. Substantial increase in the intensity of thethermal-deformation effect on the joining zonemetal in combined FW is proved also by theabsence in this zone of the SiO2 particles, whichare present in the base metal of steel AISI310and transition zone of the joints made in the«soft» mode of CFW.

Therefore, the peculiarities of structure andcomposition of metal of the AISI310 to 40Khsteel joints made by CFW and combined FW aredetermined by the character and intensity of de-formation in the contact zone at the initial, quasi-stationary and final stages of the process whenusing the above types of friction welding. Thepeculiar feature of the deformation effect in thecontact zone at the final stage of formation ofthe joints in CFW is the presence of the radialcomponent caused by applying the increasedforge pressure to the non-rotating billets. Forgingin this case provides an insignificant decrease inwidth of the transition zone and thickness of thebrittle interlayers formed at the initial stages ofthe FW process.

Combined FW with controlled deformationcreates conditions for minimisation of the proc-esses of local melting and dilution of metal vol-umes in the solid-liquid state. The effect on thejoining zone metal at the final stage of the FWprocess is characterised by the presence of bothradial and tangential components under condi-tions of the temperature gradient growing duringthe rotation deceleration process. Localisation ofdeformation and a substantial increase in its in-tensity at the final stage of welding provide dis-persion and destruction of oxide phases, as wellas displacement of interlayers outside the sectionwelded. The transition zone with a fine-grained,dynamically recrystallised structure and a dra-matic change in the concentration of alloyingelements is formed in the joint.

The data obtained served as a basis for thedevelopment of the technology for FW of thewheels of nickel superalloy Inconel 713C madeby the PIM-technology to shanks of steel 40Kh.

CONCLUSIONS

1. Characteristic features of heat-resistant steelAISI310 produced by the powder injectionmoulding technology (feed stock – «Cata-mold») are the equiaxed crystalline structurewith a grain size of about 60 μm, presence ofaustenite grains with a homogeneous and lamellar

structure, insignificant residual porosity, as wellas presence of uniformly distributed non-metallicinclusions of silicon dioxide, 2—10 μm in size,and segregations of phases with the increasedniobium content.

2. In the «soft» mode of CFW of steel AISI310to steel 40Kh, the 25 to 500 μm wide transitionzone forms within the joining zone. This transi-tion zone is non-uniform across the section, hasa lamellar structure and is an AISI310 steel based«alloy» deposited on steel 40Kh. Constant-com-position interlayers, including those correspond-ing to steel of the martensitic grade, are revealedin the transition zone on the side of steel 40Kh.Fracture of the joints in tensile tests occurs be-tween the transition zone and steel AISI310. Itis localised in segregation clusters of the redun-dant phases with the increased niobium content.

3. No segregations of low-melting point phaseswere revealed in the joints made in the «rigid»mode of CFW. However, alternate constant-com-position interlayers corresponding to steels of themartensitic and austenitic grades were fixed.These interlayers form at the initial stages of theFW process. They are not fully displaced fromthe contact zone at the subsequent heating stagesand in forging performed after stopping of rota-tion of the billets.

4. The presence of the lamellar structure con-sisting of alternate «alloys» of different compo-sitions in the joints made by CFW is indicativeof the dilution of the steels welded both in theplasticised and solid-liquid states occurring inthe contact interaction zone at the initial stagesof the FW process.

5. The technology was developed for com-bined FW with controlled deformation. Withthis technology the values of the process parame-ters are set based on the requirement to providethe certain deformation rate (burn-off rate) atthe heating stage, while the forge pressure isapplied at a stage of rotation deceleration thatis controlled following the preset program.

6. The joints made by combined FW withcontrolled deformation have fine-grained, dy-namically recrystallised structure with the up to40 μm wide transition zone that is uniform acrossthe section. The composition of metal of the tran-sition zone across its width corresponds to steelof the austenitic grade. This metal contains nopores and no segregations of alloying elementsand non-metallic inclusions. The transition zone—steel 40Kh interface is characterised by a dra-matic change in the concentration of alloying

10 9/2012

elements and absence of increased-hardness in-terlayers.

7. Parts of the materials made by the powderinjection moulding technology can be applied toprovide sound bimetal joints by using frictionwelding, subject to setting the appropriate valuesof the process parameters.

1. Dovydenkov, V.A., Krys, M.A., Fetisov, G.P.(2008) Production of metallic parts by moulding andsintering of metal-polymer compositions. Tekhnolo-giya Metallov, 6, 28—31.

2. Heaney, D. (2004) Qualification method for powderinjection molded components. P/M Sci. & Technol-ogy Briefs, 6(3), 21—27.

3. Graboj, I.E., Thom, A. (2005) Materials Catamoldof BASF Company for pressure die casting of pow-ders. Technology. Production. Application. In: Proc.of Sci.-Pract. Seminar (Joshkar-Ola, 20—21 June,2005), 71—74.

4. Salk, N. (2011) Metal injection moulding of Inconel713C for turbocharger applications. PIM Int., 5(3),61—64.

5. Froes, F.H. (2007) Advances in titanium metal injec-tion molding. Powder Metallurgy and Metal Ceram-ics, 46(56), 303—310.

6. Hamill, J.A. (1993) What are the joining processes,materials and techniques for powder metal parts?Welding J., Febr., 37—44.

7. Lebedev, V.K., Chernenko, I.A., Vill, V.I. (1987)Friction welding: Refer. Book. Leningrad: Mashino-stroenie.

8. (2006) Friction welding. Machine-building: Ency-clopedia. Vol. 1—4. Moscow: Mashinostroenie.

9. Kuchuk-Yatsenko, S.I., Zyakhor, I.V. Method offriction welding and machine for its implementation.Pat. 46460 Ukraine. Publ. 15.11.2004.

10. Son, C.-Y., Yoon, T.S., Lee, S. (2009) Correlationof microstructure with hardness, wear resistance andcorrosion resistance of powder-injection-molded speci-mens of Fe-alloy powders. Metallurg. and Mater.Transact. A, 40(5), 1110—1117.

11. Heany, D.F., Mueller, T.J., Davies, P.A. (2003) Me-chanical properties of metal injection moulded 316Lstainless steel using both prealloy and master alloytechniques. Powder Metallurgy, 1, 1—7.

12. Krug, S., Zachmann, S. (2009) Influence of sinteringconditions and furnace technology on chemical andmechanical properties of injection moulded 316L:Tech. Pap. PIM Int., 3(4), 66—70.

13. (1969) Transition joints for high temperature service.Discussion session 3. Metal Construction and BritishWelding J., 12, 134—142.

14. Gotalsky, Yu.N. (1981) Welding of dissimilar steels.Kiev: Tekhnika.

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

16. Kragelsky, I.V., Dobychin, M.N., Kombalov, V.S.(1977) Principles of friction and wear calculations.Moscow: Mashinostroenie.

17. Zyakhor, I.V. (2010) Formation of joints in frictionwelding of nickel superalloy JS3-DK to structuralsteel 40Kh. Visnyk Chernigov. DTU, 42, 148—155.

18. Smiyan, O.D. (2002) Atomic mechanism of interac-tion between ambient material and wrought metal.Fizyka ta Khimiya Tv. Tila, 4, 662—674.

19. Smiyan, O.D., Kruzhkov, A.G. (1972) About somepeculiarities of motion of diffusion gas flow in met-als. Doklady AN SSSR, 202(6), 1311—1313.

20. Smiyan, O.D., Kuchuk-Yatsenko, S.I., Kharchenko,G.K. et al. (2007) Distribution of interstitial impuri-ties within the joining zone in friction welding. ThePaton Welding J., 9, 2—5.

21. Smiyan, O.D., Zhyakhor, I.V., Nikolnikov, O.V. etal. (2008) Distribution and effect of hydrogen, oxy-gen and carbon in the joining zone of nickel superal-loy in friction welding. Visnyk Chernigov. DTU, 34,138—143.

22. Voinov, V.P., Boldyrev, R.N., Mulyukov, K.I. et al.(1976) Pulse friction welding of alloy JS6-K to steel40G. Svarochn. Proizvodstvo, 3, 28—30.

23. Zyakhor, I.V. (2000) Peculiarities of friction weldingof dissimilar metals and alloys. The Paton WeldingJ., 5, 36—44.

24. Kuchuk-Yatsenko, S.I., Zyakhor, I.V. (2002) Mecha-nism of bimetal joints formation in friction welding.Ibid., 7, 2—9.

9/2012 11

PROBLEMS OF STIRRING AND MELTINGIN EXPLOSION WELDING (ALUMINIUM—TANTALUM)

B.A. GRINBERG1, O.A. ELKINA1, A.M. PATSELOV1, A.V. INOZEMTSEV1, A.V. PLOTNIKOV1,A.Yu. VOLKOVA1, M.A. IVANOV2, V.V. RYBIN3 and Yu.B. BESSHAPOSHNIKOV4

1Institute of Metal Physics, Ural Division of the RAS, Ekaterinburg, RF2G.V. Kurdyumov Institute for Metal Physics, NASU, Kiev, Ukraine

3St. Petersburg State Polytechnic University, St. Petersburg, RF 4OJSC «Uralkhimmash», Ekaterinburg, RF

The results of investigation of structure of transition zone of aluminium—tantalum joints, having differentplane or wavy shape of the interfaces, are presented. Independently of the shape of boundaries, two typesof fragmentation were revealed: one of them is similar to fragmentation in explosion, another one – atintensive plastic deformation. The same types of fragmentation were earlier observed also in joints ofcopper—tantalum having no mutual solubility unlike the aluminium—tantalum. Melting was observed intransition zone of joints being investigated. Structure of melted areas depends greatly on the availabilityof mutual solubility in metals being welded.

Keywo r d s : explosion welding, fragmentation, transi-tion zone, local melting, mutual solubility, intermetallic

Among the whole variety of materials and weld-ing conditions the basic problem is stirring in thetransition zone near the interface [1, 2]. It is thestructure of transition zone that defines the pos-sibility of adhesion of both materials. Stirringoccurs as a result of strong external effect sup-posing large plastic deformation (including pres-sure, shear components, torque moments ofstresses, deformation non-uniformity, etc.), fric-tion of surfaces, influence of cumulative jet andother factors. However it is still unclear howstirring can occur at such a great external effectin a short time of welding procedure.

The explosion welding was carried out byCRSI of Structural Materials «Prometey» (St.Petersburg), Volgograd State Technical Univer-sity, Ural Plant of Chemical Machine Building(Ekaterinburg). Depending on welding condi-tions different joints were produced having wavyor plane shape of interfaces.

Originally the welded joints of metal—inter-metallic with normal mutual solubility were in-vestigated [3—7]. In the capacity of metal thecommercially pure titanium was selected, and asthe intermetallic – orthorhombic aluminide oftitanium (further named aluminide forbriefness). Depending on welding conditions thedifferent joints were produced which were calledfor convenience in different ways: Aw, Ap, Bw,Bp, where the low index designates the shape of

an interface (plane, wavy). For joints A the or-thorhombic alloy was used, containing 16 at.%Nb, for joints B – 23.5 at.% Nb. The moltenareas of any of mentioned joints represent genu-ine solutions where stirring occurs at the atomiclevel.

To reveal the importance of mutual solubilityof initial materials, the metals (copper—tanta-lum, iron—silver) were selected for explosionwelding which have practically no mutual solu-bility both in solid and in liquid states, moreoverthey form non-mixed suspensions in liquid state.It was found that in joints Cp, Cw of copper—tantalum [8, 9] and Dp of iron—silver [10] theareas of local melting are filled with such colloidsolutions.

Having used the obtained results the possiblereasons for good quality of a copper—tantalumjoint were revealed, due to which the explosionwelding was successfully carried out on the largeareas of plates [11] from which the chemical re-actor was designed. The body of reactor was madeof the carbon steel—copper—tantalum composite.The inner shell is composed of tantalum and onits corrosion resistance the whole structure isbased. The colloid solutions have a risk of emul-sion lamination [12]. However at the areas oflocal melting of joint Cw the lamination was notobserved [9]. But the suspension, composed oftantalum particles in the copper matrix, on thecontrary contributes to dispersion strengtheningof the joint.

© B.A. GRINBERG, O.A. ELKINA, A.M. PATSELOV, A.V. INOZEMTSEV, A.V. PLOTNIKOV, A.Yu. VOLKOVA, M.A. IVANOV, V.V. RYBIN and Yu.B. BESSHAPOSHNIKOV, 2012

12 9/2012

In the present work, which is the continuationof works [8, 9], the welded joints Ep and Ew ofaluminium—tantalum with normal mutual solu-bility are investigated. The welding parameters(angle of collision γ, speed of contact point vc)for these joints are given in Figure 1. With de-crease of welding parameters of aluminium—tan-talum (γ = 7.5°, vc = 1900 m/s) the welding didnot occur at all. The Figure shows also the weld-ing parameters for the joints Cp, Cw of copper—tantalum. As is seen from Figure 1, the parame-ters for the joints Ep and Cp, having plane bound-ary surfaces, locate near the lower boundary ofweldability. Comparing the joints Ew and Cw thewelding parameters for the joint Ew are seen con-siderably higher. Basing on the comparison ofresults obtained for the mentioned joints, theauthors tried to reveal the processes which arecommon both for the metals with mutual solu-bility as well as for the metals without it. Mean-time the simultaneous use of both joints withdifferent shape of an interface (plane, wavy) wassuccessive for the same pair of metals.

The metallographic analysis was carried outusing the optic microscope «Epiquant» equippedwith the computational complex SIAMS. Themicrostructure was investigated using transmis-sion electron microscopes JEM 200CX and CM-30 Super Twin and scanning electron microscopeQuanta 200 3D. The X-ray shooting was per-formed in the diffractometer DRON-3M in mono-chromatic Co-Kα-radiation.

The images of interfaces, obtained using scan-ning electron microscopy (SEM), plane for thejoint Ep (Figure 2, a) and wavy ones for Ew jointare given in Figure 2. For wavy interface theperiod is about 300 μm, the amplitude is 30 μm.

Fragmentation during explosion. The phe-nomenon of fragmentation as a process of split-

ting of solid body into parts (fragments) occur-ring at the intensive outer influence was knownlong ago. It came even into the name of weapon:fragmentation warhead, fragmentation shell,fragmentation bomb. The description of frag-ments separation is connected with the name ofN. Mott [13]. He together with other authors(see for example [14]) showed that using meth-ods of simple geometric fragmentation the dy-namic fragmentation of cylindrical shell can bedescribed.

We suppose that fragmentation during explo-sion investigated by Mott occurs also during ex-plosion welding. In the titanium—aluminidejoints Aw the separation of micron particles ofaluminide has a striking resemblance to separa-tion of fragments occurring during explosion butof other sizes [15]. Here it should be noted thatduring explosion the separation of fragments oc-curs in the open space, whereas in explosion weld-ing the fragments are separated at closed spacelimited by the material of plates.

For fragmentation observed in explosionwelding we used previously the notion «fragmen-tation of crushing type» (FCT). Though in both

Figure 1. Welding parameters for joints being investigated(asterisk designates the mode accepted in the work [11])

Figure 2. Interface (cross section) for the aluminium—tantalum joints: a – joint Ep; b – joint Ew

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cases the formation and separation of particlesare observed, these types of fragmentation arenot similar. FCT represents the process of split-ting into particles which either separated or abut-ted each other. Similar to explosion fragmenta-tion, FCT represents quick-running processwhich has time to occur during explosion.

Figure 3, a shows many particles of tantalumof arbitrary shape on the background of alu-

minium. We suppose that separation of tantalumparticles observed in explosion welding as wellas separation of aluminide particles mentionedabove is similar to separation of fragments inexplosion. In Figure 3, b after aluminium com-plete etching, the particles of tantalum were seenat the surface of tantalum, poured during alu-minium removal. In Figure 3, c the coarse par-ticles of tantalum are seen, reaching sizes of 20—40 μm which are much larger than those observedin the joint Ep. The possible reason for such dif-ference is that joint Ew was obtained at moreintensive outer influence than the joint Ep (seeFigure 1).

The observation of particles separation in ex-plosion welding proves the concept about frag-mentation similar to fragmentation in explosion.FCT is observed both in the metal—metal joints,as well as in metal—intermetallic ones both atthe normal solubility and also at the absence ofmutual solubility of the metals being welded in-dependently of the shape of interface. This is theevidence of dominating role of explosion in in-vestigating welding method.

FCT provides a powerful channel for dissipa-tion of induced energy, as the surface of separat-ing particles has a large total area. FCT arisesas a result of microfractures and is a process al-ternative to fracture. Instead of free surfaces,formation of which could result in fracture, theformation of surfaces due to microfractures isoccurred, which either relate to separating par-ticles or are «healed» at their consolidation. Asa result, the FCT increases the «vitality» of ma-terial preventing its fracture even at such inten-sive outer influence as explosion welding.

Fragmentation at intensive deformation. Inexplosion welding the fragmentation of one moretype is observed when the formation of new par-ticles does not take place at all. It implies frag-mentation [16], the existence of which is provedby many observations of material structure dur-ing strong deformation. Such, to the known ex-tent, traditional fragmentation [15] includespumping of dislocations (twins), formation ofnodular, cellular and banded structures, and re-crystallization.

In the present work for the joint Ep of alu-minium—tantalum the TEM-images of stronglydeformed structure of aluminium but not mixedwith tantalum were given: regions with high dis-location density (Figure, 4, a) and recrystallizedregions (Figure 4, b, c). Thus, in Figure 4, b theformation of triple butt is seen. The images ofstrongly deformed structure of tantalum are simi-

Figure 3. Microstructures of longitudinal section of transi-tion zone (a, b) and cross section (c): a – joint Ep; b –joint Ep, aluminium is etched; c – joint Ew

14 9/2012

lar to those observed for the copper—tantalumjoint Cp [15]. If to compare the images of de-formed structure of aluminium and copper [15]it is seen that recrystallization of aluminium oc-curs more completely as compared to recrystal-lization of copper. It is connected with the fact,that, firstly, aluminium is fusible and, secondly,it has higher energy of packing defect.

Though the temperature in the contact zonein explosion welding can be extremely high, how-ever at quick-running explosion effect the pro-ceeding of thermally-active processes, definingthe movement and restructuring of dislocations,is complicated and is hardly possible. It can beassumed that these processes the same as diffusionbecome possible only at residual temperaturesand stresses. These are the processes that definethe traditional fragmentation. As compared tothe traditional fragmentation, the FCT occursonly in narrow area near the interface where outerinfluence is the most intensive. However the tra-ditional fragmentation is observed somewhat far-

ther from the boundary surface. Respectively, theFCT is, by all means, the more rapid process thantraditional fragmentation. The characteristic peri-ods can be evaluated approximately as microsecondfor FCT and 108—109 μs (time of structural relaxa-tion) for traditional fragmentation.

Heterogeneities of interface surface: projec-tions, zones of local melting. In Figure 5, a theimage is composed of spots of three colors: white,black and grey. Respectively, the transition zoneis composed of areas of three types which is theresult of formation of projections at the interfacesurface. At the first time, the projections wererevealed in the titanium—aluminide joints [5].The data on chemical composition of mentionedareas forming the transition zone were obtainedusing SEM from repeated measurements. It wasshown that area of tantalum corresponds to whitecolor, aluminum area – to black, mixture ofinitial metals – to grey one.

The projections are formed most probably as aresult of diffusion-free (due to quick run of weld-ing) outburst of one metal into another (Figure 5,b). Actually, they certainly do not have perfectconic shape: the peaks are rounded and surfacesare not so smooth. Only by formation of projectionsit is possible to explain the pattern of transitionzone for joint Ep in Figure 5, a. As a result, thesurface represents chaotic relief with a large num-ber of projections of one material into another one,similarly to that observed in work [8].

Figure 4. TEM-image of structure of strongly deformedaluminium: a – cold working; b, c – recrystallized areas

Figure 5. Transition zone for joint Ep: a – microstructurefrom the area of tantalum, aluminium and zone of localmelting (longitudinal section); b – schematic image ofprojections (cones) at plane interface

9/2012 15

In the both joints Ep and Cp the drawing upof projections along some marked direction isclearly seen. It can be assumed that there is aconnection between self-organization of projec-tions and wave-formation, however it requiresspecial investigations.

If the interface was smooth, then the problemswith adhesion could arise and either reconstruc-tion of metallic links or transportation of spotdefects would be required. However presence ofprojections solves this problem: here the projec-tions play role of «wedges», linking the contact-ing materials with each other. The friction onthe surface of projection, strengthened due tothe fact that the projection itself is not smooth,contributes to adhesion of surfaces.

There is a point of view [17] that strength ofbimetal produced by explosion welding can beincreased due to preliminary profiling of surfacesof a fixed plate, which can be performed, forexample, due to longitudinal cavities and pro-jections oriented approximately parallel to thedirection of detonation spreading. A.Z. Bogunovet al. [17] consider that at the presence of pro-jections it can be spoken about explosion rivet-ing, which in case of profiling accompanies theexplosion welding.

We suppose that the similar situation occursin explosion welding due to projections formingin natural way, though their sizes are less than

in preliminary profiling. In our investigated casesthe sizes of projections are changed from mi-crometer to hundreds of micrometers, whereasafter profiling the sizes of projections varied inthe limits from several millimeters to dozens ofmillimeters [17].

Moreover, it can be assumed that existence oflower limit of weldability is connected indeedwith impossibility to form projections on theplane interface. Below the indicated limit theprojections either are not formed at all or appearto be rather small. There is obviously some criti-cal height of projections when they provide ad-hesion of surfaces. With increase in external ef-fect the projections are growing and lining up.Further, the wavy formation begins.

The X-ray diffractogram of transition zone(Figure 6, a) contains lines of tantalum and alu-minium, and does not contain lines of otherphases. The results of spot measurements ofchemical composition of grey zone obtained usingSEM showed that concentration of tantalum doesnot exceed 5 at.% here, therefore grey zone rep-resents a zone of local melting of aluminium al-loyed with tantalum. However it requires melt-ing of narrow edge of tantalum.

It is grey zone, being the zone of mixing, thatis sensitive to the fact whether materials beingwelded have mutual solubility or not. At normalsolubility they form natural solutions. At theabsence of mutual solubility the colloid solutionsarise as in case of copper—tantalum [8].

Melting along the interface. Near the wavyinterface, in contrast to plane one, the interfaceis corrugated, therefore, its longitudinal sectionrepresents a set of alternating bands of both ma-terials with almost parallel boundaries. The lon-gitudinal section of transition zone near the alu-minium—tantalum joints Ew, given in Figure 7,a, represents really a slightly inclined sectiondue to small amplitude of wavy surface. At thelongitudinal section of transition zone at the samejoint after the aluminium was etched (Figure 7,b) the bands of tantalum are seen and also par-ticles which poured after aluminium dissolution.

For comparison, Figure 8, a shows the lon-gitudinal section of transition zone for the cop-per—tantalum joint Cw. The system of bands inFigures 7, a and 8, a is not regular. If we restorethe interface according to given images, then inboth cases it will be not perfect. In Figure 8, athe numerous zones of local melting are clearlyseen having the vortex structure. As is seen fromFigure 8, b, for the both joints Cp and Cw ofcopper—tantalum the zone of local melting is

Figure 6. Diffractogram of transition zone (longitudinalsurface of section): a – joint Ep; b – joint Ew

16 9/2012

filled with colloid solution, independently of theinterface shape.

In Figure 8, a many projections of differentsizes from several micrometers to 100—150 μmare seen. Some of them begin in one band, inter-sect another one and form adhesion between thebands. The role of projections, promoting theformation of joints was mentioned above duringanalysis of the joint Ep.

As is seen from Figures 7 and 9, in the jointEw, in contrast to Cw, the zone of local melting,i.e. isolated melted areas, are not observed. In-stead of them, in Figure 9 while moving fromthe solid phase of aluminium the layers of moltenaluminium are seen, first not containing and thencontaining particles of other phase (with smalland then large density of particles). The sizes ofparticles are changed within wide range: from 50to 500 nm. Quite sharp cut of particles is par-ticularly noticeable. Further the edge of tanta-lum is followed, observed in Figure 9, a, b asilluminating line. Such illuminating boundary(«braid»), which represents a molten film, wasobserved by the authors before in the titanium—aluminide joints Ap [5].

The diffractogram taken from the longitudinalsurface of the section (see Figure 6, b) showsclearly how X-ray peaks are recorded not onlyfrom tantalum and aluminium, but also from in-termetallic phase Al3Ta. Compound Al3Ta hastetragonal crystal lattice, spatial group 14/mm,structural type – Al3Ti. Namely the particlesAl3Ta are seen in Figure 9, b. It is obvious thatmelting of tantalum is required for running ofintermetallic reaction. In fact in Figure 9, c thespherolites are seen which are the evidence oftantalum melting. The above-mentioned moltenfilm for the titanium—aluminide joint Ap has asimilar structure, where the spherolites were alsoobserved (see Figure 5, e in work [5]).

In Figure 9, d the structure of layer is givencontaining intermetallic inclusions after anneal-ing at 500 °C for 1 h. It is seen that the structureis heterogeneous but in general the coarsening ofparticles occurred as compared to those observedin Figure 9, b.

In the joint Ew the film includes edge of tan-talum of 2.0—2.5 μm width and above-mentionedmolten layer of about 40 μm width (30 μm withparticles, 10 μm – without). The observed struc-ture of transition zone is formed in the following

Figure 7. Microstructures of transition zone for joint Ew(longitudinal section): a – bands of aluminium and tan-talum; b – aluminium is completely etched

Figure 8. Microstructures of longitudinal section of theinterface for joint Cw: a – bands of copper and tantalum;numerous zones of local melting are seen; b – zone of localmelting

9/2012 17

way: at first, the melting of aluminium as morefusible one is occurred, then followed the par-tial melting of tantalum surface, formation ofsolid solution based on aluminium, and for-mation of intermetallic phase at achieving thecertain concentrations. The mentioned parti-cles are retained during dissolution of alu-minium and observed together with tantalumparticles (see Figure 7, b).

However, before melting begins in accordancewith the scenario offered above, the most quick-running process takes place: separation of solidparticles of tantalum which are seen in Figures 3,c; 7, a; 9, a.

The fact is that the molten zone in the jointEw represents a film at the interface, while inthe joint Ep the isolated zones are coordinatedwith intensification of welding condition used toproduce joint Ew as compared to the joint Ep (seeFigure 1). Here the observation of intermetallicinclusions in the joints Ew and their absence inthe joint Ep is also related to this.

Melting and adhesion. Strong explosion ef-fect during producing of the joint Ew evidenceof approaching to the upper limit of weldability.In this case the melting becomes dangerous for

integrity of the joint. However, as authors sup-pose, in favorable case (at not very high valuesof parameters of explosion effect) the melting,on the contrary, facilitates the formation of jointdue to adhesion of surfaces. In this relation, mi-crophotography is very convincing (see Figure 9,a), demonstrating transition from aluminium totantalum, occurring due to melted aluminium,aluminium with particles, and melted tantalum.It is obvious that such a transition facilitates theadhesion of aluminium and tantalum comparingwith abrupt solid-phase transition.

For other materials, namely polymers, thebest adhesive substance is solution or melt of thegiven polymer [18]. It gives a certain prove tothe possibility of adhesion also the investigatedmaterials due to their melting in explosion weld-ing, as far as during melting the problems ofwetting, adhesion and protection from contactcorrosion are instantly solved. Let us pay atten-tion to the following fact. Aluminium is oneamong few metals, the adhesion of which is im-possible without preliminary chemical treatmentof surface [19]. Only mechanical treatmentturned to be not sufficient. However aluminiumplate was not subjected to chemical treatment

Figure 9. Microstructures of the transition zone cross section for joint Ew: a – layers of melted aluminium without andwith particles; b – particles of intermetallic Al3Ta in the melt of aluminium; c – spherolites; d – after annealing

18 9/2012

before welding. Nevertheless welded joint withtantalum was produced. Here, most obviously,the remarkable feature of explosion welding wasmanifested, i.e. self-cleaning of surfaces to bewelded due to cumulative effect [1, 20].

It is known that a film of adhesive substanceshould not exceed the specified value for thick-ness [18]. As applied to welded joints it can meanthe restriction for thickness of melted area. Dur-ing intensification of welding condition and ap-proach to the upper limit of weldability the criti-cal thickness of melted area will be reached, atwhich the adhesion becomes impossible.

Thus, explosion welding comprises to a defi-nite extent also another method of joining ma-terials, namely their adhesion by formation ofmelted areas.

The electron-microscopic investigations werecarried out in the Center of collective use ofelectron microscopy of Ural Division of the RAS.

The work was carried out at financial supportof RFFI grant No. 10-02-00354, project of Pre-sidium of the RAS No. 12-P-2-1053, projects ofthe RAS UrD Nos. 12-U-2-1011 and 12-2-006-UT and State Purposeful Program of Ukraine«Nanotechnologies and nanomaterials» No.1.1.1.3-4/10-D.

1. Deribas, A.A. (1980) Physics of strengthening andexplosion welding. Novosibirsk: Nauka.

2. Lysak, V.I., Kuzmin, S.V. (2012) Lower boundary inmetal explosive welding. Evolution of ideas. J. Mat.Proc. Technol., 212, 150—156.

3. Rybin, V.V., Semenov, V.A., Sidorov, I.I. et al.(2009) Bimetallic joint of orthorhombic titanium alu-minide with titanium alloy (diffusion bonding, explo-sion welding). Voprosy Materialovedeniya, 59(3),17—31.

4. Rybin, V.V., Grinberg, B.A., Antonova, O.V. et al.(2009) Formation of vortexes in explosion welding(titanium—orthorhombic titanium aluminide). FizikaMet. i Metallovedenie, 108(4), 371—384.

5. Rybin, V.V., Grinberg, B.A., Ivanov, M.A. et al.(2010) Structure of zone of titanium joint with or-thorhombic titanium aluminide in explosion welding.Deformatsiya i Razrush. Materialov, 11, 27—33.

6. Grinberg, B.A., Ivanov, M.A., Rybin, V.V. et al.(2010) Structure of zone of titanium joint with or-thorhombic titanium aluminide in expolsion welding.II. Localized vortex melted zones. Ibid., 12, 26—34.

7. Rybin, V.V., Greenberg, B.A., Ivanov, M.A. et al.(2011) Nanostructure of vortex during explosionwelding. J. Nanosci. and Nanotechnol., 11(10),8885—8895.

8. Grinberg, B.A., Elkina, O.A., Antonova, O.V. et al.(2011) Peculiarities of formation of structure in thetransition zone of the Cu—Ta joint made by explosionwelding. The Paton Welding J., 7, 20—25.

9. Grinberg, B.A., Ivanov, M.A., Rybin, V.V. et al.(2012) Nonuniformities of interface in explosion weld-ing. Fizika Met. i Metallovedenie, 113(2), 187—200.

10. Grinberg, B.A., Ivanov, M.A., Rybin, V.V. et al.(2012) Explosion welding: Stir processes of metalswithout mutual solubility (iron—silver). Ibid.,113(11) (in publ.).

11. Frey, D., Banker, J. (2003) Recent successes in tan-talum clad pressure vessel manufacture: A new gen-eration of tantalum clad vessels. In: Proc. of Corro-sion Solutions Conf. (2003, USA), 163—169.

12. Gelfman, M.I., Kovalevich, O.V., Yustratov, V.P. etal. (2004) Colloid chemistry. St.-Petersburg: Lan.

13. Mott, N.F. (1947) Fragmentation of shell cases.Proc. Royal Soc. A, 189, Jan., 300—308.

14. Grady, D. (2006) Fragmentation of rings and shells.The legacy of N.F. Mott. Berlin, Heidelberg: Springer.

15. Grinberg, B.A., Ivanov, M.A., Rybin, V.V. et al.(2012) Fragmentation processes in explosion welding.Deformatsiya i Razrush. Materialov, 8, 2—13.

16. Rybin, V.V. (1986) High plastic deformations andfracture of metals. Moscow: Metallurgiya.

17. Bogunov, A.Z., Kuzovnikov, A.A. (2009) Productionof aluminum—steel bimetal with profiled interface.The Paton Welding J., 11, 65—68.

18. Frejdin, A.S., Turusov, R.A. (1990) Properties anddesign of adhesive bonds. Moscow: Khimiya.

19. Kapelyushnik, I.I., Mikhalev, I.I., Ejdelman, B.D.(1972) Technology of gluing of parts in aircraft con-struction. Moscow: Mashinostroenie.

20. Zakharenko, I.D. (1990) Explosion welding of me-tals. Minsk: Navuka i Tekhnika.

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EFFECT OF PALLADIUM ON STRUCTUREAND TECHNOLOGICAL PROPERTIES

OF Ag—Cu—Zn—Ni—Mn SYSTEM BRAZING FILLER ALLOYS

V.F. KHORUNOV, S.V. MAKSYMOVA and B.V. STEFANIVE.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

The effect of palladium on structure, melting ranges and technological properties of alloys of the Ag—Cu—Zn—Ni—Mn system was investigated. It was established that alloying with palladium allows improvingstrength properties of the brazed joints at an insignificant decrease of the contact angle.

Keywo r d s : brazing, diamond-hard alloy tool, cad-mium-free filler alloys, structure, melting range, contactangle, induction heating, phase transformation tempera-ture, palladium, silver filler alloy, strength of joints

Silver-base cadmium-containing filler alloys arewidely applied for brazing of various materials.Considering the negative effect of cadmium onhuman health, investigations were carried out todetermine the possibility of producing cadmium-free filler alloys with the same or similar wettingcharacteristics and melting temperatures [1—3].

As proved by analysis of literature data, cad-mium-free filler alloys of the Ag—Cu—Zn—Ni—Mnsystem proved well for brazing of diamond-hardalloy tools. Their typical representative is Ameri-can filler alloy BAg-22. This filler alloy has thefollowing chemical composition, wt.%: 48—50Ag, 15—17 Cu, 21—25 Zn, 4—5 Ni, 7—8 Mn; TS == 680 °C, TL = 699 °C, and Tbr = 699—830 °C.

The purpose of this study was to investigatetechnological properties of a standard filler alloyin brazing of hard alloy plates, determinestrength characteristics of the brazed joints andevaluate the possibility of improving the latterdue to additional alloying.

Induction heating by using a high-frequencygenerator of the VChI4-10U4 type with a dou-ble-coil inductor was employed to melt experi-mental alloys of the Ag—Cu—Zn—Ni—Mn and Ag—Cu—Zn—Ni—Mn—Pd systems under laboratoryconditions. The filler alloys were melted in alun-dum crucibles, the latter being placed in graphitecrucibles. That is, practically it was radiationheating. First the PV-200 flux was poured intoa crucible, then the Ag—Cu—Ni—Mn(Pd) chargeplaced into a nickel foil was loaded, and afterthat the flux was poured again. The charge washeated up to its complete melting and formationof the liquid pool. The ingot thus formed was

placed into a new crucible, the PV-209 flux waspoured, and the alloy was heated up to formationof the liquid pool. Zinc was added in the processof cooling the pool to a temperature of 400—500 °C. This was followed by a short-time heatingperformed several times to complete melting andachieving the homogeneous melt due to inductionand mechanical stirring.

The temperature range of melting of the alloyswas determined by differential thermal analysisusing the VDTA-8M unit in crucibles made fromzirconium oxide. Heating and cooling were per-formed in the helium atmosphere at a rate of80 °C/min. Weight of a specimen investigatedwas 1.25±0.05 g. The specimens were heated twiceto provide a good fit of the charge to the cruciblebottom and generate reliable data on thermaleffects. The thermal effects were fixed from thesecond heating curve. Solidus and liquidus of thealloy were determined from the heating curve (incooling of the alloy the substantial effect is ex-erted by overcooling prior to solidification). Atthe same time, values of the thermal effects arebest seen from the cooling curve of the alloys.

Analysis of the obtained data shows that al-loying with palladium exerts a significant effecton the phase transformation temperature andmelting range (Table 1; Figure 1, a). For in-stance, only one phase (Figure 1, b) was fixedin the Ag—Cu—Zn—Ni—Mn system alloy. The soli-dus temperature was 670 °C, and the liquidustemperature was 710 °C. In alloy with the 2 % Pdthe content the second phase in heating had anon-pronounced thermal effect, but in coolingthis effect showed up very clearly (Figure 1, c).The solidus temperature was 660 °C and theliquidus temperature was 720 °C, i.e. the solidi-fication range became somewhat wider. Alloywith the 5 % Pd content was also the two-phaseone (Figure 1, d; Table 2).

© V.F. KHORUNOV, S.V. MAKSYMOVA and B.V. STEFANIV, 2012

20 9/2012

Increase in the palladium concentration ledto increase in the solidus and liquidus tempera-tures (Figure 2).

Metallographic examinations and X-ray spec-trum microanalysis of the alloys were conductedon specimens after determination of the meltingtemperature range. For this all the alloys werecooled down to room temperature at the samecooling rate.

Alloy of the Ag—Cu—Zn—Ni—Mn system hadthree clearly defined phases (Figure 3, a). All ofthem solidified in close temperature ranges. Thatis why they had no effect on the thermal curve.It is significant that the solid solution-based pri-mary dendrites contained approximately equalamounts of copper, nickel and zinc, approxi-mately 15 % Mn, but a low amount of silver(Table 2). The silver-based (74.4 %) solid solu-tion solidified in a close temperature range. Thisphase contained no nickel, the contents of copperand manganese was several times lower, and thatof zinc was twice as low. These two phases par-ticipated in the process of solidification of the

eutectic, the typical structure of which could beseen in some regions (Table 2; Figure 3, a).

The described solidification mechanism be-came more pronounced in alloying this composi-tion with palladium. No development of solidi-fication of the primary dendrites (copper-baseddark crystals; Figure 3, b; Table 2) was detected,as the second phase (silver-based) immediatelystarted solidifying around them. Separate solidi-fication of the phases making up the eutectic took

Figure 1. Data of differential thermal analysis of the alloys:a – Ag—Cu—Zn—Ni—Mn (amorphous foil); b – Ag—Cu—Zn—Ni—Mn (cast); c – Ag—Cu—Zn—Mn—2Pd; d – Ag—Cu—Zn—Ni—Mn—5Pd

Table 1. Melting ranges of commercial filler alloy (1) and alloysunder investigation (2—4)

Alloynumber

Alloy composition TS, °С TL, °С

1 40Ag—17Cu—17Zn—26Сd 590 610

2 49Ag—16Cu—23Zn—4.5Ni—7.5Mn 670 710

3 49Ag—16Cu—23Zn—4.5Ni—7.5Mn—2Pd 660 720

4 49Ag—16Cu—23Zn—4.5Ni—7.5Mn—5Pd 670 740

Table 2. Chemical heterogeneity of the Ag—Cu—Zn—Ni—Mn(Pd) system alloys

Investigated region (phase) Ag Cu Zn Ni Mn

49Ag—16Cu—23Zn—4.5Ni—7.5Mn

Total 49.885 16.894 21.024 6.592 5.607

Dendrite (dark) 4.591 28.938 27.285 24.802 14.987

Grain (light matrix) 74.483 6.692 16.068 0 2.761

Eutectic 58.706 15.784 21.284 1.088 3.139

49Ag—16Cu—23Zn—4.5Ni—7.5Mn—2Pd

Total 64.541 11.082 16.919 1.044 3.430+(2.995Pd)

Dendrite (dark) 13.696; 11.471 43.421; 34.875 14.573; 25.190 10.905; 14.031 5.274+(12.123Pd)—7.258+(6.162Pd)

Grain (light matrix) 74.810 6.214 16.959 0 2.021+(0Pd)

Eutectic (banded) 67.617; 52.748 9.727; 20.488 15.417; 21.156 1.291; 0.335 2.564—2.219+2.850—3.057Pd

49Ag—16Cu—23Zn—4.5Ni—7.5Mn—5Pd

Total 63.603 11.371 16.723 1.932 3.239+(3.143Pd)

Dendrite (dark) 8.122 22.072 27.821 13.127 9.753+(19.038Pd)

Grain (light matrix) 70.685 7.437 15.380 0.905 2.676+(2.919Pd)

Eutectic (banded) 60.781; 59.666 13.038; 12.086 17.836; 14.758 1.105; 2.363 2.153+(5.093Pd)—2.963+(8.166Pd)

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place first. It was followed by solidification ofthe normal eutectic occupying the major part ofthe section. This mechanism was assumed to takeplace in the study by A.A. Bochvar [4]. However,it is more characteristic of organic materials, andrarely occurs in metal alloys [5]. Note that nickeland palladium were concentrated in dendrites (ofa dark colour), and were absent in the light phase(silver-based) (see Figure 3).

Increase in the palladium content of an alloycaused increase in the quantity of the primarycrystals. However, they did not grow much, asthe silver-based phase started solidifying. Theregions of separate solidification of the eutecticoccupied the major part of the section. And re-gions of the normal eutectic occurred much morerarely (Figure 3, c). Note that in this alloy the

copper-based grains contained a low amount ofnickel and palladium.

Experiments on determination of the spread-ing area over stainless steel and hard-alloy ma-terial VK8, as well as strength of the brazedjoints produced by joining these materials werecarried out by using the investigated brazingfiller alloys. The experiments were conducted onstainless steel according to GOST 23904—79 andGOST 23047—75, respectively. The results ob-tained are given in Table 3.

Structure and chemical heterogeneity of thestainless steel joints produced by using the Ag—Cu—Zn—Ni—Mn system filler alloy were investi-gated. Structure of the brazed seam turned outto be identical to that of the pure filler alloy:solid solution-based primary dendrites wereclearly revealed with the second phase and nor-mal eutectic regions solidifying around them, i.e.there was no intensive development of diffusionexchange at interfaces of the joints. This wasproved both by the data of X-ray spectrum mi-croanalysis at separate points (Figure 4; Table 4)and by the diagrams of distribution of elementsin cross sections of the seams (Figure 5).

As follows from the data of Table 4, iron andchromium, which are the main elements of the

Figure 2. Solidus (1) and liquidus (2) temperatures of thealloys depending on the palladium content

Figure 4. Location of X-ray spectrum analysis points in thebrazed seam and base metal

Figure 3. Microstructures of the experimental alloys: a – 49Ag—16Cu—23Zn—4.5Ni—7.5Mn; b – 49Ag—16Cu—23Zn—4.5Ni—7.5Mn—2Pd; c – 49Ag—16Cu—23Zn—4.5Ni—7.5Mn—5Pd

Table 3. Area of spreading and shear strength of brazed speci-mens of steel 12Kh18N10T

Filleralloy

numberAlloy system

Spreadingarea Sspr, mm

2

Average valuesof shear

strength τsh,MPa

1 Ag—Cu—Zn—Сd — 263.3 [6]

2 Ag—Cu—Zn—Ni—Mn 704.7 342.3

3 Ag—Cu—Zn—Ni—Mn—2Pd 573.0 321.3

4 Ag—Cu—Zn—Ni—Mn—5Pd 971.0 359.1

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brazed metal, were absent in the seam, the tran-sition zone was absent either. The weight contentof iron in the base metal near the interface of thejoint was 60—70 %, whereas that on the seam sidewas 0.30—0.75 %. The situation with chromiumwas similar: near the fusion boundary on the sideof the base metal its concentration was 13.5—16.5 %, and on the side of the filler alloy it was0.18—0.40 %. The seam contained no titaniumand silicon at all.

Elements of the filler alloy were present inthe seam, and were hardly revealed in the basemetal (Table 4). For example, the weight contentof silver near the interface of the joint on theseam side was within 65.0—69.5 %, and on thebase metal side – within 0—0.97 %. The similarsituation was characteristic also of copper andzinc.

The experiments on spreading were carriedout by using the 13.5 mm diameter and 4.5 mm

Figure 5. Diagrams of qualitative distribution of elements in the brazed seam metal produced by using the Ag—Cu—Zn—Ni—Mn system filler alloy

Table 4. Content of elements (wt.%) at separate points of the brazed joint

Spectrumnumber

Si Ti Cr Mn Fe Ni Cu Zn Ag

1 0.6 0.43 17.08 1.8 70.51 9.58 0 0 0

2 0.5 0.38 17.06 1.89 70.62 9.54 0 0 0

3 0.57 0.77 17.6 1.84 69.54 9.68 0 0 0

4 0.7 0.3 16.7 1.33 69.9 11.08 0 0 0

5 0.66 0.52 17.04 2.02 69.45 9.62 0.7 0 0

6 0.26 0.35 13.57 3.4 59.74 17.34 2.51 2.84 0

7 0 0 0.4 0.71 0.3 3.12 13.34 17.13 64.99

8 0 0 0.24 0.67 0.65 2.13 11.33 15.52 69.46

9 0 0 0 0.37 0.23 0.74 8.9 13.4 76.37

10 0.11 0 0 0.95 0.76 4.25 11.59 15.34 67

11 0 0 0.18 0.79 0.75 3.24 11.65 13.79 69.6

12 0.36 0.44 16.56 2.21 70.02 9.43 0 0 0.97

13 0.26 0.71 16.95 2.09 70.91 9.06 0 0 0

14 0.5 0.41 16.99 1.44 70.71 9.95 0 0 0

15 0.6 0.59 17.33 1.71 68.7 9.81 0 0 1.26

16 0.53 0.27 17.18 1.92 70.44 9.67 0 0 0

17 0.34 0.5 17.7 1.62 69.45 9.34 0 0 1.06

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thick plates of hard alloy material VK8 and filleralloys in the form of cubes measuring 4 × 4 ×× 4 mm.

Prior to brazing the specimens were degreasedwith acetone (alcohol), then the filler alloy beingstudied was placed at the specimen centre, andflux of the PV-209 grade having the reactivitytemperature range equal to 600—850 °C was de-posited on the top over the entire perimeter of aspecimen [7]. Heating of the specimens was per-formed by using high-frequency generatorVChI4-10U4 with a frequency of 440 kHz andpower of 10 kW. The double-coil inductor wasemployed for these studies. A rack of heat-resis-tant ceramics with a channel to attach a thermo-couple for measuring the temperature of heatingof a specimen up to complete spreading of thefiller alloy was installed inside the inductor.Holding for three seconds was performed after

melting of the filler alloy. After that heating wasstopped. The degree of wetting of a solid sub-strate with the filler alloys was determined byevaluating the area of spreading of a molten filleralloy (three specimens for each filler alloy) andcontact angle between the substrate and a spreaddroplet of the filler alloy metal (Figure 6). Thecontact angle was determined on the sections cutout from the specimens normal to the wettingplane (Figure 7). The data were obtained by us-ing the AutoCad software.

Specimens for mechanical tests were brazedin air by using high-frequency heating. Flux inthe form of a paste mixed up with water waspreliminarily deposited on the hard alloy plates(VK8+VK8, VK15+VK15 etc.). Then a filler al-loy charge was placed between the hard alloyplates, and heating was performed up to completemelting of the flux and filler alloy, as well as upto formation of the joint.

Shear strength was determined by the V.N.Bakul Institute for Superhard Materials of theNAS of Ukraine by using a special attachment(to the tensile testing machine P-05). As seen

Figure 6. Appearances of specimens 1—3 for investigationof spreading of filler alloys 1—4 (a—d) over hard alloy ma-terial VK8

Table 5. Results of tests of the VK8 joints brazed with filler al-loys of different systems

Filleralloy

numberFiller alloy system

Spreading area

Sspr,* mm2

Contactangle θ,

deg

Shearstrengthτsh, MPa

1 Ag—Cu—Zn—Cd 95.30 17 245

2 Ag—Cu—Zn—Ni—Mn 136.32 10 225

3 Ag—Cu—Zn—Ni—Mn—2Pd 111.06 16 250

4 Ag—Cu—Zn—Ni—Mn—5Pd 112.01 13 315

*Spreading area was determined by using high-frequency heatingfollowing a non-standard procedure

Figure 7. Cross section of specimens after spreading tests using filler alloys 1—4 (a—d) (see Table 5)

24 9/2012

from the data given in Table 5, alloying of theexperimental filler alloy with palladium substan-tially increased strength of the joints.

Filler alloys of the investigated systems wereemployed to manufacture a batch of bits for drill-ing the wells to produce dispersed methane,which increased several times the length of deep-ening of the wells [8].

Therefore, application of the environmentallyclean filler alloys of the Ag—Cu—Zn—Ni—Mn andAg—Cu—Zn—Ni—Mn—Pd systems instead of filleralloys of the Ag—Cu—Zn—Cd system for brazingof rock and metal cutting tools makes it possiblenot only to avoid the harmful effect of cadmiumon human organism, but also to provide the in-creased-strength brazed joints.

1. Roberts, P.M. (1978) Recent developments in cad-mium-free silver brazing alloys. Welding J., 10, 23—30.

2. Timmins, P.F. (1994) The development of Ag-basedbrazing alloys. Ibid., 10, 31—33.

3. Khorunov, V.F., Maksymova, S.V., Stefaniv, B.V.(2010) Effect of tin additions on structure and tech-nological properties of brazing filler metals of theAg—Cu—Zn system. The Paton Welding J., 7, 16—21.

4. Bochvar, A.A. (1956) Metals science. Moscow: Me-tallurgizdat.

5. Taran, Yu.N., Mazur, V.I. (1978) Structure of eutec-tic alloys. Moscow: Metallurgiya.

6. Petrunin, I.E., Lotsmanov, S.N., Nikolaev, G.F.(1973) Brazing of metals. Moscow: Metallurgiya.

7. Klochko, N.A. (1981) Fundamentals of the technol-ogy for brazing and heat treatment of hardalloytool. Moscow: Metallurgiya.

8. Khorunov, V.F., Maksymova, S.V., Stefaniv, B.V.(2010) Manufacture of drill bits for production ofdispersed methane in mine working. The PatonWelding J., 6, 41—43.

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SUBSTANTIATION OF «LEAK-BEFORE-BREAK»CRITERION FOR VERTICAL CYLINDRICAL TANKS

FOR OIL STORAGE

V.V. KNYSH, A.Yu. BARVINKO, Yu.P. BARVINKO and A.N. YASHNIKE.O. Paton Electric Welding Institute, NASU, Kiev, Ukraine

The paper gives the results of experimental studies of cyclic crack resistance of samples of butt-weldedjoints on 06GB-390 steel in the case of oil storage tanks at stable development of a through-thicknessfatigue crack, initiating from the notch surface along the fusion line, from the moment of its initiation upto reaching the length of about 30 mm. Large-scale samples of welded joints of 650 × 160 × 20 mm dimensionswere tested at harmonic alternating zero-to-tension stress cycle with σmax = 2/3σy. It is shown thatapplication of rolled sheets of the above-mentioned steel for design rings of tank wall allows, at applicationof «leak-before-break» criterion, eliminating extended fractures in the wall welded joints.

Keywo r d s : fatigue crack, development of through-thickness fatigue crack, «leak-before-break» criterion, cy-clic loading

Over the recent decades, because of instabilityof oil price in the world oil market, many coun-tries are building tanks of 50,000—100,000 m3

and greater capacity to create the required stockof this product. Requirements to ecological safetyof such constructions are becoming more strin-gent with every year, accordingly. Normative re-quirements as to steel mechanical properties, inparticular in Z-direction, for rolled sheets of tankwalls and their welded joints have grown signifi-cantly. However, in view of various peculiarities,defects in the wall welded joints cannot be elimi-nated completely, which at wall operation undercyclic loading conditions can become the sourceof initiation and development of tough fatiguefractures. Under such conditions, the problem ofestablishing the fatigue life of the stage of stablesubcritical propagation of through-thickness fa-tigue crack in design rings of tank wall is becom-ing particularly urgent. Appearance of a darkspot on the white wall from seeping oil from theformed through-thickness crack, can be a reliablecriterion of crack detection in those cases whenthe time of its growing up to critical size willallow guaranteed detection of the crack underthe conditions of technical examination of tanksaccording to the applied procedures. Such a cri-terion of crack detection in pressure vessels wascalled «leak-before-break» [1].

The purpose of the proposed work was experi-mental substantiation of the possibility of appli-cation of «leak-before-break» criterion to vertical

tanks for oil storage, in which the lower ringsare made of a new class of high-quality steel06GB-390.

Application of new generation niobium-con-taining steels 06GB-390, 06G2B-440 [2] with im-pact toughness values KCV—40 ≥ 120 J/cm2 8to 40 mm thick (furtheron strength classes willbe omitted in the work) and with fine-grained iso-tropic structure for design lower rings of the wallsessentially improved their crack resistance. In thisrespect study [3] is of interest, which shows thatstabilization of critical crack opening displacementwithin δcr = 0.18—0.20 mm is observed for a numberof tested steels with KCV—20 ≥ 80 J/cm2. In viewof that, it is logical to expect that at applicationof 06GB steel for lower rings of tank wall andpossible formation of a fatigue crack in it, itsgrowth rate will give enough time for detectionof the above crack at subcritical stage of its de-velopment.

As was noted above, at application of «leak-before-break» criterion under the conditions oftank operation, it is important that the durationof the stage of stable crack growth in the linear(second) section of fatigue fracture diagram [4],expressed by the number of tank filling—emptyingcycles, allowed enough time to enable detectionof a through-thickness crack after oil outflowingfrom it, and taking the required safety measures.In this work flat samples with a transverse buttweld were tested to substantiate the criterion«leak-before-break» for a wall of oil storage tank(Figure 1). In samples made from rolled sheetsof 06GB steel, cross-section with removed weldreinforcement was equal to 160 × 19.5 mm, andlength was 650 mm. Maximum stress of harmonicloading cycle was taken, allowing for maximum

© V.V. KNYSH, A.Yu. BARVINKO, Yu.P. BARVINKO and A.N. YASHNIK, 2012

26 9/2012

value of calculated hoop stresses in the tank wall[5] found from the following condition:

σh ≤ Ryγw/γn,

where Ry = 350 MPa is the calculated resistance;γw = 0.8 is the coefficient of operating conditionsof wall rings; γn = 1.1 is the fitness-for-purposefactor.

In view of the above, maximum stress of al-ternating loading cycle was taken as σmax == 260 MPa. Such a level of hoop stresses allowsfor tolerances [5] for deviation of the geometricalshape of new tank wall from the design value.Considering that the tank wall is a thin-walledshell of up to 60 m diameter and up to 20 mheight, after 15—20 years of service the norms[6] allow exceeding the initial geometrical de-viations (foundation sagging under the wall, ap-pearance of dents, etc.) 2 times. It is natural thatunder these conditions, in individual regions ofthe wall hoop stresses will be significantly higherthan the initial values, that was allowed for bycoefficient γw = 0.8, but these stresses should notexceed the calculated values for this steel con-sidering the periods of technical examination oftanks. Growing of the initial surface notch in thesample at the first stage up to formation of athrough-thickness crack, and then growth of thiscrack leads to reduction of sample net sectionarea and increase of stresses in the sample thatto a certain extent simulates the noted local in-crease of hoop stresses in the actual tank wall.

Initial crack-like surface notch was made insample center at a distance from the weld fusionline with specially ground cutter. Length of

semielliptical notch 2a was 38 mm, and its depthwas 6 mm at notch rounding-off radius of 25 mm.All the samples were tested in hydro-pulsatorTsDM 200pu at uniaxial zero-to-tension alternat-ing stretching with 5 Hz frequency. Duration ofthe stage of surface crack initiation from thenotch and its growing through sample thicknessup to appearance of through-thickness crack isgiven in [7].

Fatigue testing of samples was conducted atroom temperature; however, considering thatrolled sheets of 06GB steel of 8—50 mm thicknessmaintain their high plastic properties in the tem-perature range T = +20 — —40 °C [2], it can beassumed that the obtained test results are correctup to temperature T = —40 °C. During testingthe initial length of the formed through-thicknesscrack was measured on sample side opposite tothe notch and number of cycles, at which itformed, was measured. Furtheron cycle numberwas recorded at its elongation by each 3—6 mm.At crack length 2a ≈ 40 mm its tip developedplastic contractions and testing was stopped. Re-sults of fatigue testing of the samples are givenin the Table. Figure 2 gives the dependence offatigue crack length on number of alternatingloading cycles.

Table data show that at crack length 2a ≤≤ 30 mm in sample net section stresses did notexceed the calculated resistance of 06GB steel(Ry = 350 MPa). After the crack has grown tothe length of 2a > 40 mm in sample net section,maximum stresses reached values close to steelyield point, i.e. appearance of observed plasticcontractions at such a crack length is not related

Figure 1. Schematic and dimensions of samples for testing welded joints of 06GB steel for cyclic crack resistance: 1 –notch in the fusion zone; 2 – weld reinforcement was removed from two sides

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to limit state of cyclic crack resistance, but isdue to limited dimensions of the sample. Suchhigh stresses will not develop in the rings ofactual tank wall at specified crack length, be-cause of many times smaller weakening of thewall by the considered defect.

Application of empirical dependence betweenK1C and KCV specified in the norms allows de-termination of critical crack length 2acr for thegiven welded joint at minimum temperature oftank service T = —40 °C. In keeping with [8]

K1C = 0.1√⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯0.1 E1 — μ2 KCV—40 , MPa⋅√⎯⎯m. AtKCV—40 = 246 J/cm

2 we get K1C == 236 MPa⋅√⎯⎯m for the weld. At possible in-creased hoop stresses in the wall σh = Ry == 350 MPa, proceeding from the known relation-ship for through-thickness crack of critical length2acr, K1C = σh√⎯⎯⎯⎯πacr, we get value acr = 145 mmfor the given σh and 2acr = 290 mm, respectively.For new tanks of 50,000 m3 capacity in keepingwith the norms [5], hoop stresses in the wall are