Micro Structural Evolution and Mechanical Properties of Nb-Ti Micro Alloyed Pipeline Steel

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Available online at www.sciencedirect.com_-@ ScienceDirectJOURNAL OF IRON AND STEEL RESEARCH, INTERNATIONAL. 2011, 18(2): 57-63

Microstructural Evolution and Mechanical Properties ofNb-Ti Microalloyed Pipeline Steel

LAN Liang-yun, QIU Chun-lin, ZH AO De-wen, GA O Xiu-hua(State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang 110004, Liaoning, China)

Abstract: Th e correlation between microstructure s and mechanical properties of a Nb-Ti microalloyed pipeline steelwas investigated. Th e results revealed that with decreasing the finish rolling tempera ture and th e cooling sto p tem-peratu re, the matrix microst ructure was changed from quasi-polygonal fe rrite to acicular ferrite, as a result of im-provement of both st ren gth and low tempera ture toughness. By means of electron backscattered diffraction observa-tion, an effective acicular ferrite packet contained several low angle boundaries o r subboundaries plates which madeimporta nt contributions to improvement of stre ngth. It was found that many fine quasi-polygonal f errite grains withhigh angle boundaries as the toughening structure were introduced into the acicular ferrite matrix to refine effectivegrain size and improve the toughness.Key words: pipeline steel; acicular ferrite ; mechanical property; electron backscattered diffraction

T o improve the transmission efficiency, pipelinesteels must have high str eng th, high toughnes s, andexcellent corrosion Recently manystudies have been carried out by optimizing chemicalcomposition and thermomechanical control process-ing (TMCP) in order to obtain excellent mechanicalproper ties of pipeline stee lsC2-67. Whe n pipelinesteels are subjected t o different rolling processes,the micro structures obtained usually consist of po-lygonal ferrite (PF) , quasi-polygonal ferrite (QF)or massive ferrite, granular bainitic ferrite (G F) andbainitic or acicular ferrite (BF or AF) in terms offerrite microstructures classified by Krauss and otherr e s e a r c h e r ~ ~ ~ - ~ I .he A F , first described in the early1 9 7 0 ~ ~ ~ ,as been well known as the optimum mi-cro stru ctur e with an excellent combination of highstrength and good low temperature toughness inweld steels . Recently many researche s have beendone on the A F formed in pipeline steels, because itsmorphology and properties were t o some extent sim-ilar to the AF formed in weld steelsC0-163.However,the morphological feature and crystallographic orien-tation of such kind of microstructure formed in pipe-line steels are still controversial issues according to

Ref. [ l l ] , Ref. [15] to Ref. [18].In this work, different multiphase microstruc-

tures, such as AF together with QF and dispersedmartensite/austenite (M/A) constituent, were ob-tained by applying different rolling processes. Thepurpose of this stu dy is comparative analysis of t heeffects of d ifferent multiphase microstructures onmechanical properties of an API X70 pipeline steeland confirmation of the AF microstructure crystallo-graph ic characteri stics by means of e lectron back-scattered diffraction (EBSD) and scanning electronmicroscope (SEM).1 Materials and Experimental Procedure

Chemical composition for the experimental steelis listed in Table 1. In order to obtain low cost andhigh st reng th of hot rolled plate, the amount of carbonelement was controlled slightly high but much lower

Table 1 Chemical composition of the experimental steel(mass percent, % I

C Mn S P Si A1 Nb Ti0.12 1.13 0 . 0 0 3 0.017 0.25 0.035 0.028 0.019

Foundation Item: Ite m Sponsored by National Natural Science Foundation of China (51074052) ; Fundamental Research Funds for CentralBiography:LAN Liang-yun(l983- ), Male, Doctor Candidate; E-mail: lly. 1iangyunBgmail. com; Received Date: November 3 , 2009

Universities of China (N100607001)


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than the maximum value prescribed in the API spec5L standardC6'. Th e Nb and Ti elements as main mi-croalloyed elements were added in it. In view of theelements of S and P great damage to toughness,these contents (i n mass percent) were limited to0. 0 0 3 % and 0. 017% respectively.

Th e materials, 100 mm X 100 mmX 150 mm insize, for hot rolling were cut from the forged slabs.Th e hot rolling experim ent was carried out throughvarying TM CP conditions on pilot rolling mill withtwin rolls of 45 0 mm in diameter. Details of rollingconditions are shown in Table 2. All steels were hotrolled to 1 0 mm-thick plates through a two-stagecontrolled rolling process. The firs t stage rollingwas controlled in the recrystallized austenite region

at high temperature w ith heavy reduction in order torefine the original austenite size, while the secondstage rolling w as controlled in non-recrystallizedaustenite in low temp erature region above A r3 .

Specimens for tensile and impact t ests were cu tfrom middle of th e rolled plates in the transver sal di-rect ion. Round tensile specimens with a gauge diameterof 8 mm and a gauge length of 40 mm were tested atambient temperature at a crosshead speed of 5 mm/minusing a SANS 10 kN servo-hydraulic machine. To eval-uate the Charpy impact ene rgy, sub-size (7. 5 mm X10 mm X 10 mm ) standa rd Charpy V-notch speci-mens were conducted over a temperature range from-60 "C to room temperature in accordance with thestandard method of ASTM-E 8M.

Table 2 The rolling schedule and measured processing temperaturesof the experimental steelsRough rolling region Finish rolling region Acceleration cooling

Steel Finishing Beginning Finishing Cooling stop Cooling rate/temperature/% temperature/% temperature/C temperature/% ( % * s-1)

A 1040 940 925 680 33B 1 0 4 0 940 840 675 22C 1060 870 820 590 32D 1 0 6 0 868 810 560 34

-Schedule/mm A-D 100--73+50+33 33+25+19.5+15.5+12+10

The longitudinal direction plates of th e hotrolled steels were polished and etched by 2 % of Nitalsolution, and microstructures were observed by FEIQuanta 600-scanning electron microscope ( SEM ).After SEM observation all the specimens were elec-trochemically polished in a solution containing 70 %of ethanol, 22% of distilled water and 8% of per-chloric acid (in volume percent), operating at 30 Vand -25 OC. EBSD analyses were carried out in SEMequipped with a conventional EBSD system at theresolution step of 0.2 pm on the area of 70 pmX80 pm.T o quantify the effective grain size ( d ) betweenhigh-angle boundaries (>15") defined by M D Fuen-tes et alCln1,misorientation profiles were obtainedfrom lines traced on orientation imaging maps, and alinear intercept method was used to compute thisvalue. In order to examine cleavage facet size whichis closely related to the toughness of material, cleav-age fracture surfac es of the C VN specimens fracturedat -60 'C were a lso observed by SEM.2 Results and Discussion2 .1 MicrostructuresI t is a valid way to achieve excellent mechanical

proper ties of pipeline ste els by optimizing microstruc-tures subjected to different rolling Inthis experiment, microstructures are composed ofmany different phases including QF, P (pearlite) ,GF and AF or BF. Th e microstructures of th e A andB stee ls rolled at high finish rolling and cooling stoptemperature are shown in Fig. 1 (a) and ( b > respec-tively. Th e microstructur es are mainly composed of

(a ) A steel; (b ) B steel; (c) C steel; (d) D steel.Fig. 1 SEM micrographsof all the steels


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Issue 2 Microstructural Evolution and Mechanical Properties of Nb-Ti Microalloyed Pipeline Steel 59

QF, together with P and little AF. The QF grainshave irregular and jagged boundaries with high dislo-cation density and little dispersed M/A constitu-entC7. Since the QF transformation can be accom-plished by short-diffusion of interstitial (C) and sub-stitutional ( N b ) atom across transformation inter-facesC8, it can be observed th at carbides discontinu-ously appear at the QF grain boundaries. The micro-structures of the C and D steels are shown in Fig. 1 (c )and (d) respectively. A s low finish rolling and cool-ing stop temperature, the A F formed in middle tem-perature is the major microstructure together with GFand QF. The AF with excellent mechanical p ropert iesmay be an optimal microstructure for pipeline steels

because i t includes high dislocation density and dis-persed M /A constituent betwe en the f errite plates.2.2 Tensile properties

Th e tensile properties of all the steels tested atroom temperature are listed in Table 3 , as well astheir stress-strain curves are represented in Fig. 2.Th e strength of AF together with G F and QF micro-stru ctur es is much higher tha n that of QF togetherwith P and A F microstructures. By decreasing thefinish rolling and cooling stop temperature the sub-stru ctur e and dislocation with the high density areformed and retained in austenite, which increase thenucleation sites of the AF and promote the AF transfor

Table 3 Tensile properties and impact absorbed energy of all the steelsAbsorbed energy/J

-1 0 C -4 0 C -6 0 CSteel Yield s t rength/M Pa Tensi le s t rength /MPa Elongat ion/ % Yield ratio/ %A 485 595 23 0. 82 75 55 47B 49 0 600 20 0. 82 77 65 40C 545 670 25 0. 81 135 145 130D 58 0 700 20 0. 83 147 160 11 4

0 0.02 0.06 0.10 0.14 0.18True strain

Fig. 2 True strain-true stress curves of all experimental steelsmation, and it is beneficial to refine the microstruc-ture and improve the ~t re ng th ~ ~ . eanwhile theyield rati os, the ratio of yield str ength to tensilestrength, maintain a low constant about 0. 82 ratherthan increase with th e increase of yield strength .Th at is to say, the microstructures of A F togetherwith GF and QF can improve the yield stre ngth andkeep the low yield ratio. Thi s is consistent with theresult of Kim studiedC7.

2.3 Impact toughness and fracture behaviorsThe CV N impact toughness of all the s teels tested

at various temperatures are also listed in Table 3.Th e toughness of the A and B.steels is very low at

all test temperatures, while the C and D steels hav-ing A F matrix perform excellent low temperaturetoughness. Fig. 3 ( a ) and ( b ) show SEM fracto-graphs of CVN specimens of the A and D steels frac-tured at -6 0 OC respectively. Cleavage facet size, a nimporta nt fac tor influencing the toughness of material,can be measured by fractogra phs, a nd it is related tobut appears to be somewhat larger than the effectivegrain size which will be discussed below since cleav-age cracks are deflected at the high misorientationb o u n d a r i e ~ ~ ~ - ~ ~T h e cleavage facet size of the Asteel is measured to be about 10 pm , larger than thatof the D steel (about 6 pm). On the other hand, thefract ions of cleavage frac ture area of the A and Bsteels are larger than that of the C and D steels ac-cording to macrofractograghs. Especially no cleavagefractur e phenomenon appears on th e C steel fracto-graph fractured at - 0 C, which illustrates that theC steel has the best excellent toughness in this ex-periment. According to Eqn. (1 ) derived for the duc-tile-brittle transition temperatures T D B T of low-alloyferritic steels established by F B P i ~ k e r i n g ~ ~ ,heT D B T of mate ria l can be improved by 22 C when con-taining P of 1 0% in microstructures. There fore itcan be inferred that the P causes a great damage tothe toughness of A and B steels.

T~~~=-19+44zela+700wNtn 11. 5d-+2. 2 ~ p(1)


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( a ) A steel; ( b ) D steel.Fig. 3 SEM fractographs of CV N specimens fractured at -60 c

where, wN(f)s nitrogen not combined as a stable ni-tride; and d is effective grain size.2 .4 EBSD analyses

It is well known tha t good toughness of the mi-crostructures is closely related to small effectivegrain size or high density of high angle gra in bound-aries which can be measured by EBSD technique.High angle boundaries can act as obstacles to cleavagepropagation, forcing the cleavage crack to change

the microscopic plane of propagation. Low angleboundaries seem to have no influence on the tough-ness of steels. From this reason, in order to discussfracture mechanics, it is more convenient to use theconcept of crystallographic packet, defined by A FGourgues et alc13*191,orresponding to grains or setsof adjacent units sharing the same crystallographicorientation.

Fig. 4 (a ) to ( d ) represent image quality mapsand orien tation image maps of the A and D steels. The

( a ) , (b ) A steel; ( c ) . (d ) D steel.Fig. 4 Image quality maps and orientation image maps


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Issue 2 Microstructural Evolution and Mechanical Properties of Nb-Ti Microalloyed Pipeline Steel 6 1

image quality maps a re built up by grain orientationcontrast. Owing to the overlapping of the Kikuchipattern coming from the small size of A F plates ana-lyzed in t he conventional EBSD system, the shape ofA F plates could not be accurately represented in theFig. 4 (c). Th e orientation image maps display grainorientations in unique colors dedicated to each crys-tallographic plane index, and each point that belongsto these planes exhibits with it s neighbor points amisorientation lower than a certain ( 15") toleranceangle. Some dispersed dark gray colors correspond-ing to the low pattern quality are usually caused bylath boundaries or carbide particles, even pearlitephase s hown in orientation image map of the A steel.

Fig. 5 ( a ) t o ( d ) reveal misoriention maps ofthe A to D steels respectively after cleaning up thelow patt ern points. Light grey lines represent lowangle boundaries (10"- 5") , while dark grey linesrepresent high angle boundaries (15" or more). Ac-cording to t he misorientation maps of the A to Dsteels shown in Fig. 5 , the fraction of high angleboundaries bf th e A to D steels is 87.7%, 86.5%,89. 7% , and 87. 8% respectively. High angle bounda-ries play an important part in the toughness of mate-rial because they are usually thought as a sufficient

(a ) A steel; (b ) B steel; (c ) C steel; (d ) D steel.Fig. 5 Misorientation maps of all steels

condition for the crack arrest which can force thecleavage clacks to change the microscopic plane ofp ropaga t i~n"~- '~ ]In term s of t he linear interceptmethod, the effective grain size of th e A to D steelsis 6. 6 , 6. 9 , 4. 9 , and 5. 8 pm respectively. Thesmall effective grain size is the primary determinantattributio n to excellent toughne ss of material. In aword, these results can properly explain that thetoughnes s of the C and D steels are more excellentthan that of th e A and B steels, because the crackpropagation path in AF matrix steel is much morebent than that in QF matrix steelCzo1.Comparing theorientation image maps and the misorientation mapswith the SEM microstructures, it is interesting ob-served that each QF grain become effective grain be-cause of its high angle boundaries with misorienta-tion of 15" or more , which can be confirmed byFig. 6 (b ). An A F grain size shown in the imagequality figures does not become an effective grainsize, but the AF packet containing several adjacentgrains with low misorientation becomes the effectivegrain, which can be confirmed by Fig. 6 (a). And itis worth pointing out that each A F crystallographicpacket contains several morphological packets formed byparallel plates observed by SEM. Therefore theSEM images define much finer microstructures thanthose visualized in the orientation image maps.Misorientation profiles of different microstructureswere investigated along lines selected in orientationimage maps of t he D ste el, as shown in Fig. 6 (a) to(c). Fig. 6 ( a ) re presen ts the misorientation anglevaried w ithin an A F packet along the line ( a b ) . I t isrevealed that the point to point and point to originmisor ientat ion angles are lower than 7. 5" and 15" re-spectively. And the curves exhibit a fluctuant varia-tion which indicates that the A F microstructuresconsist of high intricate plates with lots of internallow angle boundar ies and high density of dislocation.It is a major factor to result in the improvement ofthe strengt h of t he C and D steels. Meanwhile thelength of this AF packet is more than 20 pm, which in-dicates that the size of A F packets to some extend enl-

0 4 8 1 2 1 6 2 0

I W

60

200 4 8 12 16

LU

151050

Pointtoorigin

1 3 5 7 9

Fig. 6 Misorientation profiles of different microstructures corresponding to different lines selected in Fig. 5 (d)


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arge instead of refining the effective grain size. Th eA F plates in pipeline steels, for this feature , are dif-ferent from the A F plates in weld steels where eachAF plate is considered as the effective grain. T hereason is tha t the nucleation sites of the AF in pipe-line stee ls are not so many as that in weld steels ac-cording to relevant literatureC'0"4-'7~21-2z1 Fig. 6 (b )represents the misorientation angle varied along theline ( c d ) across several QF grains. The misorienta-tion angle abruptly boosts up to 60 " when encounte-ring QF grains boundaries and the distance betweentwo adjacent high angle boundaries is about 2 . 5 -4 pm. Therefore QF as the toughening structure inth e C and D steels makes a significant contributionto refine the effective grain size. The misorientionangle varied along the line (ef) selected in a GFgrain is shown in Fig. 6 (c). A few low angle bounda-ries in the G F grains are much less than those in th eAF packets. T he cleavage crack propagation in mul-tiphase microstructures of A F together with QF isschematically described by Fig. 7c223 .From the frac-ture viewpoint , he A F crystallographic packets andfine QF grains act the effective grains and theirboundaries can effectively deflect the cleavage crackpropagation resulting in excellent low temperaturetoughness.

QF

Fig. 7 Schematic illustration of cleavage crackpropagation in microstructures AF and QF

3 ConclusionsIn this experiment, the pipeline steels with dif-

ferent microstructures were produced by varying th ehot rolling conditions and the effects of different mi-cros truc ture s on mechanical p ropert ies were investi-gated by mechanical properties tests and EBSD analysis.

1) The steels rolled at low finish rolling tem-perature composed of A F and G F together with fineQF presented excellent mechanical properties ; bu tthose rolled at high finish rolling temperature mainlycomposed of QF together with P and little AFshowed low impact absorbed energy values since ir-regularly dispersed P and carbides in QF grainboundaries caused the low temperature toughness.

2 ) The fine QF grains as the toughening struc-ture in A F matrix were t he effective grains with highangle boundaries, while an A F packet containingseveral similar crystallographic orientations platescould become the effective grain. Th is kind of AF incrystallographic properties was different from theAF formed in weld ste els because of the differentnucleation mechanisms.References:c11

C 2l

C3l

c 41

C5l

[ e lc71

C8l

c91

C l O l

Clll

c151

C161

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Micro Structural Evolution and Mechanical Properties of Nb-Ti Micro Alloyed Pipeline Steel