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NEELACHAL ISPAT NIGAM LTD

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STUDY ON TEMPERING OF Niobium –Vanadium Containing H.S.L.ASteels

Date:30.04.2018

Name: ARPITA HAZRA

Email ID: [email protected]

K-Bank

mailto:[email protected]


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Confidentiality StatementThe information in the document mentioned is not confidential and have been taken references from

various sources as specified.

AbstractIn this paper the effect of Tempering on Nb-V containing H.S.L.A steel properties ,i.e. strength&toughness studied. In an earlier work a study was done to find out the optimum slab-reheating temperature of this steel to have optimum properties.

The reheating temperatures were varied between 1100 0 -1250 0 c. Reheated samples weregiven three types of cooling rates, i.e . water quenching, interrupted quenching and air cooling.

The mechanical properties of samples were studied by tensile testing and Vickers hardness.The observation w .r .t the properties were justified through optical metallography. In theprevious study samples (reheated at 11500 c) were found to be most optimum with respect tostrength and toughness considered.

The present study is a continuation of the above study in the direction of enhancing toughnessparameter without any detriment to the strength by tempering at different temperature andtime scale. Properties of samples were studied by Vicker’s hardness, Microhardness andtensile testing along with study on optical micrograph reveal the possibility of phenomena likesoftening, secondary hardening, refinement of elongated(Pan-caked) ferritic structure. Fromobservation of the above phenomena with respect to time and temperature scale it hasbecome possible to derive optimum treatment condition for maximum toughness advantages.

About the Author

Arpita Hazra . Mgr (PPC),NINL. She was associated with NINL for last 16 years.


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Intended Readers

Introduction1. The advanced technology demands new variety of engineering materials for the differentstructural applications. The development of H.S.L.A (High strength low alloy) steels can beviewed as the opening of a new horizon in metallurgy. The key properties which make itsuperior to plain carbon steels are as following:-

1) High yield strength .2) Good toughness property.3) Low impact transition temperature value.4) Good formability property5) Good weldability property.6) Fatigue and abrasion resistance properties

To attain optimization between different properties thermomechanical controlled rollingoperation is the best way. Variation in re heating temperature helps in different mechanicalproperties. The constant increasing demands made on H.S.L.A steels require increasinglyhigher strength levels with improved toughness . This has led to the development of thismaterials in quenched and tempered form in addition to their application in the conventionalnormalized and as rolled condition.

The micro alloying elements which are found to be very effective in imparting strength andtoughness includes Nb, V, Ti, Al,N2, Mo. The mechanisms of strengthening can be attributedto

1) Strengthening through ferrite grain refinement2) Solid solution strengthening by the solute micro alloying elements3) Precipitation hardening and dispersion strengthening4) Texture hardening.5) Dislocation strengthening.6) Fibre strengthening7) Substructure strength


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Description2. LITERATURE REVIEW

Introduction to Microalloyed Steel:

High strength low alloy steels having elements Nb,V, Al or Boron as micro-alloying additions areknown Microalloyed steels. The total concentration of alloying elements is less than 0.3% as individualor in combination.

Properties :

1. Tensile Property

Microalloyed steels have a very fine ferrite grain size, in the range of ASTM 12-14. The finest grainsize is obtained by controlled rolling. The yield strength increases with decreasing grain size. Inaddition ,the fine dispersion of alloy carbides results in precipitation hardening. Typical contributionsfrom various strengthening mechanisms to the total yield strength are as follows:

Strengthening mechanism Contribution, MPa

Iron Lattice 35

Solid solution 135

Precipitation 45

Dislocation 50

Grain size 260

Total 525

It is evident that the major contribution comes from grain refinement. It has the additional advantagein that the impact transition temperature(ITT) decreases with decreasing grain size, where as itincreases with other strengthening methods. The increase in the strength of micro alloyed steelsresults in lower elongation values.


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2. Formability:

Formability can be defined as the ability of a metal to be formed without necking or tearing during acold-forming operations such as drawing, stretching or bending. Micro alloyed steels have poorerformability than the conventional low carbon steels.

3. Weldability:

The excellent weldability of mild steel remains unaffected in micro-alloyed steel.

Thermo mechanical Treatment:

The development of high strength materials, coupled with good formability, has always been the aimof material scientists. The demand for such a combination has increased considerably during the lastfew decades because of the ever increasing applications in space, deep sea and high pressuretechnology. Alloying, mechanical working, heat treatment, grain size control, nuclear irradiations aresome of the techniques which may be taken recourse to individually or in combination for this purpose.Individually ,each one of them has a limited strengthening effect. In order to attain enhancedproperties, various combination of these unit operations may be adopted advantageously. Thermomechanical treatment is one such combination. As the name suggests, it is a combination of heattreatment and mechanical working. However, this statement does not define the term thermomechanical treatment in the strict metallurgical sense since any combination of heat treatment andmechanical working cannot be termed a thermo mechanical treatment. The term thermo mechanicaltreatment refers to that treatment in which plastic deformation is carried out in such a way that phasetransformation is affected by it.

The principle behind thermo mechanical treatment is that plastic deformation results in the productionof various crystal defects such as vacancies, dislocations, sub-grain boundaries and stacking faults.These defects severely affect the phase transformation in metals and alloys by providing nucleationsites and aiding diffusion process. These in turn affect the kinetics of phase transformation andmorphology of phase(s) formed. Therefore, an important aspect of thermo mechanical treatment isthat phase transformation should occur under conditions of increase lattice defects. This means thatthere should be overlapping of two processes, namely ,heat treatment and mechanicaldeformation .Both processes may either proceed simultaneously or with a time gap.

Objectives of tempering:-

Hardening treatment develops maximum hardness, excellent wear resistance and high strength levelsin the steel. At the same time, it affects adversely properties such as, ductility, toughness and impactstrength. It also imparts brittleness to steel because of the internal stresses developed by quenching.The extent of brittleness depends on the chemical composition and cooling rate. The degree ofbrittleness depends on the chemical composition and severity of cooling rate. Thus, steel in ashardened state is unsuitable for some service conditions and rarely used as such. A relatively stablestate can be attained by providing thermal energy to the steel. This results in decrease in internalstresses and reduction in the degree of brittleness. Such a process which consists of heating


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hardened steel below the lower critical temperature , followed by cooling in air or at any other desiredrate, is known as tempering.

Tempering treatment lowers hardness, strength and wear resistance of the hardened steel marginally.However, this marginal loss is adequately compensated by relieving of internal stresses, restoration ofductility and toughness and transformation of retained austenite. The higher the temperingtemperature, the more is the restored ductility and tougher the steel, However due attention should bepaid towards hardness, strength and wear resistance properties. This is the aim of the hardeningtreatment. Proper tempering treatment results in optimum combination of mechanical properties.Elastic properties of steel are also affected by this treatment. Hardening follower by tempering is theonly conventional heat treatment process suitable for improving elastic limit of the steel.

General aspects of carbide & nitride precipitates in steels.

The main alloying elements Nb,V & Ti all form stable carbides & nitrides which possesses the FCCstructure. Solubility products have been determined for the relevant carbides and nitrides in austeniteas a function of temperature is shown below:-

Sufficient solubility of alloying elements in austenite is essential at the reheat temperature to allow fineprecipitates to occur during rolling passes at lower temperature. The main refinement fromprecipitation for grain growth control is achieved during rolling as the temperature progressively fallsand fine carbo nitrides precipitate formed. These precipitates increase the strain required forrecrystallisation at a given temperature and restrict the movement of recrystallised grain boundaries &


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thus increase the strength.

Quenched and Tempered steels

The transformation product obtained by rapid cooling on low carbon steel except thin sheet is nonmartensitic. The lower the finish rolling temperature, the better are the properties due to the finer prioraustenite grain size which give rise to a very fine acicular ferrite structure. This treatment alsointroduce an extra amount of defect structure which leads also the better properties. This latter is veryeffective in increasing the strength when the final rolling occurs at temperature where the austeniteferrite transformation takes place in very low carbon steels. However, the impact properties may beimpaired.

This structure can be improved further with grain refinement technique by usual additions or bycontrolled rolling coupled with dispersion hardening by Nb or V. Some work on Nb, V containingferrite-pearlite steels had suggested that these materials might be responsive to quenching heattreatment which would give fine grain size(for adequate toughness) while subsequent tempering wouldlead to strengthening the precipitation of Nb or V carbide. The secondary hardening effect obtainingby the alloy carbide is well established. Thus there is an obvious need to evaluate the effect ofquenching and subsequent tempering treatment on the properties of low carbon Nb,V containing microalloyed HSLA steels.

As expected, an increase in C or N content brought about an increase in the as-quenched hardness.Acicular ferrite structure has a high strength by virtue of its fine grain size and high dislocation density,coupled with good impact toughness. So we can go for lower carbon content without impairing itsstrength, comparing with other ferrite-pearlite steels. The presence of Mn increase the hardenability sothat the acicular ferrite is produced even on air cooling.

During tempering the major role is played by the Nb,V. Niobium,Vanadium not only promote andmaintain the finer prior austenite grain size to achieve the fine grained acicular ferrite structure butalso allow some precipitation strengthening during tempering. For the precipitation of NbC or V4C3

relatively higher tempering temperature or time is required. The acicular ferrite structure contain a lotof dislocations and dark field examination revealed very fine precipitates on dislocations. Thehardness on tempering in the range of 5000C -6000C range is due to alloy carbides( or carbo nitrides)precipitation on dislocation. Over aging also deteriorates the hardness value above 6000C .

There was no yield point in the as-quenched material or in the lightly tempered material. The yieldpoint occurred on tempering at 5500C for 4 hrs.

The as quenched hardness varied according to the (C+N) content of the steel though the Mn has asignificant effect. The tempered peak hardness are mainly influenced by the (V+N) content. The asquenched and tempered strength & hardness are considerably greater than those achieved in othersteels of this group. The higher alloy content would be expected a sufficient effect on thetransformation characteristics to account for the superior strength and inferior toughness of the steel.

The hardenability of the low carbon Vanadium containing steels is very low. The ferrite transformation


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region would be very close to the time axis and transformation would commence at a relatively hightemperature(7000C -8000C ).Also rapid cooling would increase the velocity of the transformationreaction and nucleation would be favoured relative to growth. This would explain why the morphologyof the ferrite in the quenched material are acicular ,rather than the more usual polygonal grainobserved after slow cooling.

The effect of increasing Mn and (C+N) content was to lower the temperature of transformation.Therefore ,on quenching ,this steel showed a substantially a smaller amount of free acicular ferrite andthe structure was mainly bainitic. Thus the high strength and poor impact properties can be achievedon the basis of the difference in structure.

With respect to tempering treatment acicular ferrite is different from martenste. The latter owing mostof its strength to carbon in solution. Acicular ferrite owes its properties to a very fine ferrite grain sizeand a high concentration of mobile dislocations which give rise to good toughness and ductility.Precipitation of, for example ,NbC by tempering strength low carbon acicular ferrite but reduces itstoughness . Precipitation of carbon from solution in martensite by tempering decreases the strengthbut usually increase the toughness.

The high dislocation density desired in the structure of the acicular ferrite must play an important rolein raising the strength of the material. Since a large proportion of these dislocations were not ‘pinned’they would be capable of restricted movement under an applied stress.

In most of the steels the improvement in strength obtained on quenching were accompanied by asubstantial lowering of the transition temperature. It is observed that presence of pearlite and/or alloycarbides (certain distributions only) have a deleterious effect on the ductile/brittle transitiontemperature of low carbon steels. The absence of both these phases in the quenched structure couldaccount for improved transition temperature. Thus the metallurgy and design philosophy of acicularferritic steels should not compare with the standard quenched and tempered low alloy engineeringsteels.

The size and spacing of the precipitates would tend to lock the dislocations and a high applied stresswould be necessary to free them. The locking of dislocations would result in a high yield strength anda high Y.S/UTS ratio.

Though the precipitation of vanadium carbide down to temperatures below 5000C has been observedthe precipitation of niobium carbides has not been observed at tempering temperature below6500C.But this phenomenon also depends on the tempering time and the quenched structure.Obviously, it is much easier for the NbC to nucleate on dislocation sites and this could well account forthe precipitation of NbC at temperature as low as 5000C.

In steels containing low Nb ( or V) with low C the transition temperature was not affected aftersecondary hardening. Presumably this is because the volume fraction of the precipitates is quite small.When larger amount of these elements are present the transition temperature deteriorated ontempering. This results suggest that in these low carbon steels it is possible to achieve improvedstrength( by carbide precipitates during tempering) without prejudicing the impact properties of the


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material. By a judicious combination of composition and mechanical and thermal treatment it might bepossible to achieve a precipitate (and dislocation) distribution which would give increased strengthwithout sacrifice in toughness.

3. EXPERIMENTAL WORK

3.1 Alloy Design

With the present understanding of physical metallurgy of HSLA steels it has become imperative torestrict percentage of carbon within 0.08%,Niobium within 0.025% and Vanadium within 0.1% and Tiwithin 0.025% in order to have a fine austenite grain size, as well as to achieve better effect onprecipitation and dislocation hardening.

The chemistry of the experimental was selected very judiciously. The chemistry of the alloy isshown in table 4.1

3.2 Thermo Mechanical Controlled Process Design

Vanadium Carbide dissolves in austenite completely around 11000C . TiC & TiN is soluble inaustenite at high temperature than vanadium carbide. So the selected alloy was soaked at fourtemperature i.e.11000C,11500C,12000 C& 12500C.

3.3 Experimental Heat Making

The experimental steel making has been carried out in R & D centre of M/s TISCO LTD. in thefollowing manners:

Low carbon scrap in the form of small cut pieces was melted in a 50 kg induction furnace. Toavoid oxidation loss during melting, protective slag cover of lime was maintained throughout themelting operation. This slag was flushed off after complete melt down when the temperature of theliquid melt was in the range of 15500C to 15700C.

In order to achieve low active oxygen (less than 5 ppm) Aluminium deoxidation was done byadding pre calculated amount of aluminium shots, followed by addition of Fe- Si and electrolytemanganese ferro-alloys like Fe-V, Fe-Nb & Fe-Ti. All these additions were performed within a veryshort span of time of 2-3 minutes. To protect the steel from re-oxidation ,liquid steel was covered withsynthetic slag which was a mixture of different calcium bearing compounds like finely powdered freshlime (CaO) ,Fluorspars (CaF2) and small quantity of calcium carbide(CaC2) etc. The liquid steel wastapped at a temperature of 16000C in the form of two ingots of 25 kg each. The cast iron ingot mouldswere preheated for half an hour and teeming was done by top pouring using a hot top. The steelchemical composition was analyzed by taking samples while tapping.

3.4 Forging


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The cast ingots were forged at TISCO. The dimensions of the forged bars were 200 mm x70mm x30mm.

3.5 Expermental Rolling

Twelve pieces were cut from the forged bar. Dimensions of each bar was 70 x 30 x 25 mm.These twelve pieces divided into four groups (A,B,C& D).Samples of A,B,C & D were soaked at11000C(T1),11500C(T2) ,12000C and 12500C(T4) respectively. Then the samples were rolled using 10H.P rolling mill. Total reduction was given 60 percent to each sample. Total number of pass for rollinga sample was eight and the minimum reduction maintained in the first three passes was 40%.Afterrolling three samples of each group were given three different cooling rates i.e. aircooling(AC),interrupted quenching(IC) and water quenching (WC).

3.6 Tempering Treatment

To observe the tempering effect of the H .S.L.A steel selected for present study, a particular groupof samples i.e. group ‘B’ (soaked at 11500C and then air cooled, interrupted quenched and watercooled after rolling) was selected. The tempering temperatures were 5000C, 5500C and 6000C and theperiod of tempering times were 3 hrs, 4 hrs,5 hrs and 6hrs.

3.7 Tensile and Charpy Impact properties

Tensile Test

Tensile samples of 6 mm x 6mm dimensions were prepared and carried out in 10 MT Instron tensiletesting machine using 25 mm extensometer, cross head speed of 1.25 mm per minute.

Impact Testing

Impact testing has been carried out at three temperatures i.e. 460C, 00c and room temperature. Thespecification of the samples used for impact testing were 7.5 x 10 x 55 mm and 2 mm v notch.

3.8 Hardness Measurement

Vicker’s hardness of all the heat treated samples has been measured using combine Rockwell-Brinell-Vicker’s hardness tester. Micro-hardness of all the tempered samples has been measuredusing Polyver-Met microhardness tester.

3.9 Metallographic studies

Samples were cut from the heat treated samples and were ground, polished and etched with 2%Nital reagent. We took the microphotograph of some selected tempered samples (5500C).

4. RESULTS

The outcome of this work in terms of mechanical properties, hardness values are shown below.


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Schematic diagram for rolling schedule has been shown. The photomicrograph are shown in Figs a,band 1-8.

Table 4.1

Compositions of the Micro-Alloyed Steel

%C %Mn %S %Si %Al(T) %Al(Soln) % Nb %V %Ti

0.05 1.65 0.020 0.297 0.053 0.05 0.035 0.11 0.0016

Table 4.2

Rolling Schedule

Pass no 1 2 3 4 5 6 7 8

StockThickness(mm) 25 21.5 18.5 16 14 12.5 11.25 10.25 10.00

% of reduction 14 14 13.5 12.5 11 10 8 2.5

Table 4.3

Sampleidentification

Soakingtemp

Coolingmodes

Tensile PropertiesVicker’s

Hardnessvalues

CHARPY ImpactProperties(Joules)

U.T.S(MPa)

Y,S(MPa) Y.S/UTS

Ratio

%Strain

atBreak

-46 0

C00C RT(250C)

Ai 11000C

Air Cooling(A1) 216 33 59 76

Interruptedquenching 581.7 395.7 0.68 24.5 218 22 38 73


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Waterquenching 887.5 684.5 0.77 20.76 329 24 37 53

Bi 11500C

Air Cooling

(B1)599.3 505.4 0.84 36.60 212.2 26 61 76

Interruptedquenching 767.9 648.1 0.84 32.16 305 28 57 67

Waterquenching 794.3 671.2 0.845 33.56 305 23 36 43

Ci 12000C

Air Cooling(C1) 573 422 0.736 46.76 201.8 25 80 96

Interruptedquenching 620.2 509.8 0.821 40.16 219.9 47 100

Waterquenching 656.1 517.3 0.786 41.96 239 29 43

Di 12500C

Air Cooling(D1) 592.4 437.3 0.738 42.48 226 6 26 41

Interruptedquenching 620.3 495.5 0.79 40.72 216.4 36 35 47

Waterquenching 790 683.1 0.86 35.44 278 27 40 67

Table 4.4

SoakingTemp

Coolingmodes

TemperingTemp(0C)

TemperingTime(hrs)

Tensile Properties

U.T.S( Mpa) Y,S (Mpa)

Strain atmax.load

11500C Interruptedquenching 550

4 639.3 1.577 0.2360

5 634.7 545.600 0.2524


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6 725.6 1.318 0.1948

11500C Waterquenching 600

4 759.9 1.189 0.1908

5 726.5 658.000 0.1780

6 730.7 1.021 0.2104

11500C Interruptedquenching 600

4 682.4 265.100 0.1904

5 675.6 4.644 0.2244

6 704.9 332.800 0.1532

11500C Waterquenching 550

5 682.1 2.027 0.2216

6 614.9 296.000 0.2300

Table 4.5

Micro Hardness values of Bi samples after Tempering

Sample identification Tempering Temp(0C) Duration ofTempering (hr.)

Micro-Hardnessvalues

B2 i.e soaked at11500

C and theninterrupted quenched

after rolling.

500 3 128.20

4 101.3

5 145.7

6 131.5

550 3 111.4

4 132.4

5 134.0

6 139.8

600 3 100.6

4 51.8


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5 135.2

6 118.8

Sample identification Tempering Temp(0C) Duration ofTempering (hr.)

Micro-Hardnessvalues

B1 i.e soaked at11500

C and then air cooledafter rolling

500

3 88.0

4 97.40

5 111.7

6 93.5

550

3 100.1

4 108.0

5 108.0

6

600

3 116.7

4 127.4

5 117.1

6 127.1

Sample identification Tempering Temp(0C) Duration of Tempering(hr.)

Micro-Hardness values

B3 i.e soaked at11500 Cand then water quenched

after rolling

500 3 60.0

4 62.4

5 132.7

6 105.1

550 3 165.3

4 138.4


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5 133.7

6 116.9

600 3 98.0

4 106.5

5 102.4

6

Table 4.6

Vicker’s Hardness values of Bi samples after Tempering


Vicker’s Hardness value

B i.e soaked at11500 Cand then air cooled after

rolling

500

3 209

4 227.1

5 212

6 220.8

550

3 247

4 240

5 213

6 213

600

3 202.6

4 199.5

5 100.1

6 236.5




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B i.e soaked at11500 Cand then interrupted

quenched after rolling

500

3 230.0

4 288.0

5 230.8

6 234.6

550

3 221.0

4 233.6

5 23.50

600

3 216.0

4 208.9

5 208.10

6



B i.e soaked at11500 Cand then water quenched

after rolling

500 3 247.0

4 260.0

5 269.0

6

550 3 229.2

4 276.0

5 269.0

6 271.0

600 3 246.0

4 230.8

5 216.4

6


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5. DISCUSSION

In our preliminary work with respect to optimization of homogenization temperature yielded that for Nb-V kind of steel 11500C is the best choice with respect to soaking temperature and this was supportedby study on strength and ductility. The samples of the same material was taken for the present studyand tempering was performed at temperature 5000C,5500C and 6000C for 3 hrs,4 hours, 5hrs and 6hrs. Heated samples have been studied with respect to micro-structure , microhardness and tensileproperties.

Discussion on microstructure

Microstructure of parent material with air cooling shows coarse ferritic grain with considerablethickening at boundaries having occasional bainite regions at triple points. From earlier studies thiskind of structure has not been recommended for high toughness purpose, with interrupted quenching(Fig. 4,5) apolygonal ferritic grains with bainitic region have been observed ,grain boundary thickeningis least due to diffusion restriction at lower temperature. With water quenching samples (Fig.6,7,8)reveal existence of fine elongated ferritic grains (nearly acicular) along with apolygonal ferritic grains.From the study of previous rolling schedules and related finish rolling temperature for 11500C andselected pass design in present study, finish rolling temperature is considered to come belowsolutionizing temperature of Nb and V . Recrystallization is thus inhibited due to presence of serrationand bulges at boundaries and sub-boundaries, those arise from deformation and act as the potentialsites for ferrite nucleation.

On tempering at 5000C in air cooled samples carbide precipitates at grain boundaries and showsboundary thickening with increasing temperatures. In interrupted quenching and water quenchedsamples carbides precipitated at the boundary of apolygonal grains.

At 6000C considerable thickening of boundaries have been observed without any other remarkablechanges. At high temperature existence of extensive coalescence has been observed. The abovestructures are common with tempering and have not been shown. Some special features have beenobserved in samples treated at 5500C (Fig.1-8), except for air cooled samples, the interruptedquenched and water quenched samples show evidence of lateral precipitation at elongated ferritegrains. Along this lateral precipitation elongated finer grains show shrink at different regions of


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boundary. This may be precipitation of carbides at serrations and buldges of elongated ferrite grainboundaries and subsequent diffusion of solute atoms to the precipitate resulting in drag of grainboundaries and formation of very fine ferritic grains. At higher temperature(6000C) and higher durationof holding time (5 hrs and 6hrs) further coalescence of carbides and ferritic grains results due tomigration of boundaries.

Observation on mechanical properties

From study of micro-hardness data shown in table 4.5 samples tempered at 5500C shows best kind ofresults for interrupted quenched samples at 5 hrs and water quenched samples at 4 hrs. This is duefine microstructure obtained fine microstructure obtained by lateral refining of elongated ferrite grains.

Tensile data shows best combination of strength & toughness for untempered samples for interruptedquenching. With subsequent tempering at specified temperatures and duration shows that sampleswith interrupted quenching for 4 hrs & 5 hrs. gives the best combination of strength & ductility that canbe achieved with composition mentioned in table 4.1.

With the above observations it is realized from the present study that the benefit of partial solutionizingof alloying elements with proper selection of homogenizing temperature & proper selection of finishrolling temperature may produce a microstructure between acicular and polygonal ferrite of elongatedmorphology with interrupted quenching . Subsequent tempering at proper temperature and durationmay give rise a structure with lateral refining of ferrite grain that present best combination of strengthand ductility. Proper selection of time and temperature is essential because further heating forgreater duration will give rise coalescence of carbide and subsequent migration of ferrite boundary togive rise coarsening effect.

6. Conclusion:-

1. Nb, V steel is an unique combination to obtain strengthening effect contributed by grainrefining, precipitation strengthening and solution strengthening.

2. Interrupted quenching gives rise compromise between strength and ductility without anytempering.


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3. Tempering of Nb, V steel gives rise initial softening, refinement of elongated morphology bylateral precipitation and secondary hardening effect in the range of 500-600 0C and timeinterval 3 hrs to 6 hrs sequencial order .

4. Lateral precipitation and subsequent shrinkage in the region of serration and bulges havebeen observed as a determining phenomenon in this kind of steel for improving strength andtoughness.

5. A schedule for heat treatment for this material can be suggested as following:

i) Heating at 11500 C

ii) Soak for 2 hrs

iii) Roll upto a temperature just below solutionizing temperature

iv)Water quenching upto room temperature or upto 5500C and subsequent air-cooling

v) Tempering at temperature near 5500C for 4 to 5 hrs


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5 . DISCUSSION

In our preliminary work with respect to optimization of homogenization temperature yielded thatfor Nb-V kind of steel 11500C is the best choice with respect to soaking temperature and thiswas supported by study on strength and ductility. The samples of the same material was takenfor the present study and tempering was performed at temperature 5000C,5500C and 6000C for3 hrs,4 hours, 5hrs and 6 hrs. Heated samples have been studied with respect to micro-structure , microhardness and tensile properties.

Discussion on microstructure

Microstructure of parent material with air cooling shows coarse ferritic grain with considerablethickening at boundaries having occasional bainite regions at triple points. From earlier studiesthis kind of structure has not been recommended for high toughness purpose, with interruptedquenching (Fig. 4,5) apolygonal ferritic grains with bainitic region have been observed ,grainboundary thickening is least due to diffusion restriction at lower temperature. With waterquenching samples (Fig.6,7,8) reveal existence of fine elongated ferritic grains (nearly acicular)along with apolygonal ferritic grains. From the study of previous rolling schedules and relatedfinish rolling temperature for 11500C and selected pass design in present study, finish rollingtemperature is considered to come below solutionizing temperature of Nb and V .Recrystallization is thus inhibited due to presence of serration and bulges at boundaries andsub-boundaries, those arise from deformation and act as the potential sites for ferrite nucleation.

On tempering at 5000C in air cooled samples carbide precipitates at grain boundaries andshows boundary thickening with increasing temperatures. In interrupted quenching and waterquenched samples carbides precipitated at the boundary of apolygonal grains.

At 6000C considerable thickening of boundaries have been observed without any otherremarkable changes. At high temperature existence of extensive coalescence has beenobserved. The above structures are common with tempering and have not been shown. Somespecial features have been observed in samples treated at 5500C (Fig.1-8), except for aircooled samples, the interrupted quenched and water quenched samples show evidence oflateral precipitation at elongated ferrite grains. Along this lateral precipitation elongated finergrains show shrink at different regions of boundary. This may be precipitation of carbides atserrations and buldges of elongated ferrite grain boundaries and subsequent diffusion of soluteatoms to the precipitate resulting in drag of grain boundaries and formation of very fine ferriticgrains. At higher temperature(6000C) and higher duration of holding time (5 hrs and 6hrs)further coalescence of carbides and ferritic grains results due to migration of boundaries.


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Observation on mechanical properties

From study of micro-hardness data shown in table 4.5 samples tempered at 5500C shows bestkind of results for interrupted quenched samples at 5 hrs and water quenched samples at 4 hrs.This is due fine microstructure obtained fine microstructure obtained by lateral refining ofelongated ferrite grains.

Tensile data shows best combination of strength & toughness for untempered samples forinterrupted quenching. With subsequent tempering at specified temperatures and durationshows that samples with interrupted quenching for 4 hrs & 5 hrs. gives the best combination ofstrength & ductility that can be achieved with composition mentioned in table 4.1.

With the above observations it is realized from the present study that the benefit of partialsolutionizing of alloying elements with proper selection of homogenizing temperature & properselection of finish rolling temperature may produce a microstructure between acicular andpolygonal ferrite of elongated morphology with interrupted quenching . Subsequent temperingat proper temperature and duration may give rise a structure with lateral refining of ferrite grainthat present best combination of strength and ductility. Proper selection of time andtemperature is essential because further heating for greater duration will give rise coalescenceof carbide and subsequent migration of ferrite boundary to give rise coarsening effect.

6 Conclusion:-

1 Nb, V steel is an unique combination to obtain strengthening effect contributed by grainrefining, precipitation strengthening and solution strengthening.

2 Interrupted quenching gives rise compromise between strength and ductility without anytempering.

3 Tempering of Nb, V steel gives rise initial softening, refinement of elongatedmorphology by lateral precipitation and secondary hardening effect in the range of 500-600 0C and time interval 3 hrs to 6 hrs sequencial order .

4 Lateral precipitation and subsequent shrinkage in the region of serration and bulgeshave been observed as a determining phenomenon in this kind of steel for improvingstrength and toughness.

5 A schedule for heat treatment for this material can be suggested as following:

i) Heating at 11500 C

ii) Soak for 2 hrs

iii) Roll upto a temperature just below solutionizing temperature


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iv)Water quenching upto room temperature or upto 5500C and subsequent air-cooling

v) Tempering at temperature near 5500C for 4 to 5 hrs

Acknowledgement

References1.Physical Metallurgy and design of Steels, F.B Pickering

2.Heat Treatment,R.Sharma.Sharma.

3.Processing,Microstructure and properties of H.S.L.A steels, edited by Dr. A.J.Deardo

(i) Carbide precipitation in HSLA sreel, R.W.K Honeycombe

(ii) Mice alloying and thermo mechanical treatment, Michale Korchynsky

(iii) Microstructure and mechanical properties of comparable HSLA controlledrolled ,microalloy, C-Mn Steels, L. Roberson Link and R.W Armstrong

(iv) . Precipitation and recrystallization during TMP of complex microalloyedHSLA steels,A.C Kneissl, G.Posch,C.I.Garcia,A.J.Dardo.

4. High strength low alloy steels, International Iron and Steel Institute Committee ontechnology,

Brussels 1987.


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5. J.M.Gray, D. Webster and J.M.Woodhead, 1965 ,precipitation in mild steelscontaining small

Additions of Niobium. journal of the iron and steel institute, vol.203,p.818.

6. J.J.Irani, 1966, “Quenched and Tempered low carbon steels containing Niobium orVanadium” . Journal of the Iron and Steel Institute, vol.204 p-702-710

7. Noami Matsumura and Mosahru Tokiazane,1986,” Austenite grain refinement andsuper

plasticity in Niobium micro alloyed steel “ Transaction of the iron and steel instituteof Japan,vol,p-315-321.


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