Effect of dynamic strain aging on mechanical

Effect of dynamic strain aging on mechanical properties of vanadium microalloyed steel S. Gu¨ndu¨z and R. C. Cochrane In the present work, the dynamic strain aging behaviour in microalloyed steels has been examined using a C – Mn – Al – V – N steel, which occasionally exhibits low toughness in the subcritical heat affected zone (HAZ). This may be attributed to dynamic strain aging, whereby materials show lower ductility and higher yield strength owing to the interaction between mobile dislocations and diffusing solute atoms. The research has shown that the high temperature tensile behaviour of C – Mn – Al – V – N steel depends on the presence of clusters believed to be of vanadium and carbon atoms. The interaction between dislocations and clusters of vanadium and carbon atoms at 200 – 450uC changes the work hardening rate and contributes to dynamic strain aging as confirmed in the present work. These interactions may also decrease toughness in the subcritical HAZ and lead to the subcritical embrittlement observed in the C – Mn – Al – V – N microalloyed steel. MST/5502 Dr Gu¨ndu¨z ([email protected]) is in the Karabu¨k Technical Education Faculty, Zonguldak Karaelmas University, 78200, Karabu¨k, Turkey and Professor Cochrane ([email protected]) is in the School of Process, Environmental and Materials Engineering, Department of Materials, University of Leeds, Leeds LS2 9JT, UK. Manuscript received 30 April 2002; accepted 2 July 2002. # 2003 IoM Communications Ltd. Published by Maney for the Institute of Materials, Minerals and Mining.

Introduction When carbon and nitrogen atoms diffuse towards dislocations in the stress field of the dislocations they form atmospheres which lock the mobile dislocations This immobility of the dislocations then results in a higher strength and a lower ductility in steels. The atmospheres can form either after prestraining and unloading, or during straining. Depending upon whether the straining and aging processes take place sequentially or simultaneously, strain aging can be classified into two types: static strain aging and dynamic strain aging.1,2 Dynamic strain aging and static strain aging have different effects on the mechanical properties of materials. Generally, dynamic strain aging affects mainly the plastic deformation or work hardening behaviour, while static strain aging affects the yield strength of materials. The focus of many studies of static strain aging has been the reappearance of the lower yield point upon aging after prestraining and unloading, and this phenomenon is considered to be a universal manifestation of static strain aging.3 – 8 Microalloyed steels typically contain niobium, titanium or vanadium, either singly or in combination, and their specific effects may be influenced by other alloying additions such as aluminium, boron or indeed any of the other more conventional alloying elements used in steel manufacture. The effects of the microalloying elements are also strongly influenced by thermal and thermomechanical treatments.9 It is common to assume that when microalloying additions are made the static strain aging response is reduced, owing to a much reduced interstitial content as a result of the formation of microalloy carbides and nitrides. The dynamic strain aging (DSA) response is much less documented, but most aluminium treated steels show lower increments in flow stress at high temperatures than those of silicon killed or balanced steels. Niobium treated steels display an intermediate level of DSA response. Under some processing conditions, i.e. normalising, particularly in vanadium treated steels the ferrite can remain supersaturated in the microalloying element such that a classical age hardening response can occur.10 In these circumstances it is reasonable to suppose that there would be interactions between carbon and the microalloying element, similar DOI 10.1179/026708303225001984

in nature to those assumed to occur during low temperature creep. Glen11 postulated that precipitation occurs on dislocations produced during creep deformation. Furthermore, carbon (or nitrogen) could be trapped by the substitutional element, as shown by Baird and Jamieson,12 at temperatures at which precipitation would be expected to have ceased, leading to dynamic strain aging effects. One consequence of this would be for the strain hardening rate to increase, ultimately reducing ductility. Studies of toughness changes adjacent to welds occasionally show a deterioration in Charpy impact energy in the subcritical region of the weld heat affected zone (HAZ).13 While such changes do not have a major structural significance, they can result in failures in meeting specifications, such as those encountered in shipbuilding or offshore oil platform fabrication. These zones of lower toughness occur in regions corresponding to temperatures of approximately 200 – 500uC, and, when noted, have been interpreted as evidence of strain aging.14,15 However, such zones are not pronounced after normalising, and this could imply that the interstitial content differs between rolled and normalised steels. A more radical interpretation is that dynamic strain aging is more pronounced or occurs over a wider temperature range. Such a hypothesis is supported by the observation of Glen and others that the strain aging and temperature dependence of yield in aluminium treated normalised steels is consistent with a very low interstitial content. The present investigation was undertaken to determine whether dynamic strain aging can occur in a C – Mn – Al – V – N microalloyed steels under conditions relevant to weld HAZs. A considerable amount of information exists about the transformation characteristics and precipitation behaviour in this steel, which indicates that a range of precipitate sizes and distributions can be produced using particular heat treatments.16

Experimental procedure The steel chosen to study strain aging behaviour in simulated HAZ conditions was a commercial high strength C – Mn – Al – V – N steel. The chemical composition of the Materials Science and Technology

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2 Gu¨ndu¨z and Cochrane Dynamic strain aging of V microalloyed steel

conventional mechanical properties were derived from the foregoing information: lower yield stress (LYS), ultimate tensile stress (UTS), elongation to fracture (%), work hardening index n, and initial work hardening rate. In the case of continuous curves, yield strengths were determined from load – extension curves by the 0.2% strain offset method. To determine the work hardening index n it was necessary to calculate true stress s and true strain e and log true stress and strain for the region of uniform plastic deformation up to maximum load. The slope of the latter relationship defines the work hardening index n, given17 by s~ken

:

:

:

:

:

:

:

:

:

:

:

:

:

:

: (1)

log s~log kzn log e where k is a constant. Grain sizes were measured using intercepts along a test line oriented at 45u to the rolling direction. At least 500 grain boundaries were counted for each specimen. After counting grain boundaries, the total length of the intersecting lines was recorded and the following formula was used to determine mean linear intercept grain size of the ferrite18 ia ~‘(1{fp )=na

1 Cutting plan of ‘I beam’ steel

steel was (wt-%) Fe – 0.14C – 0.40Si – 1.41Mn – 0.02P – 0.03S – 0 .03Cr – 0.005Mo – 0.05Ni – 0.035Al – 0.017N – 0.005Ti – 0.10V. Steel was supplied in the form of a column 3176317 mm by 307 mm. The flange thickness was y25 mm. The steel column was cut according to the diagram shown in Fig. 1. Flanges 1 and 2 were cut from the web and then into strips A – D. These were subsequently cut into smaller strips A1 – D2. From each of these numbered strips, five specimens were obtained. All specimens except those in the as received condition were austenitised before testing, by heating to 900uC, and holding for 1 h at temperature. A batch of specimens were then continuously cooled in air (AC), while others were continuously cooled in a stainless steel cylinder (SSC), to retain variable amounts of vanadium as well as carbon and/or nitrogen in solid solution. After austenitisation at 900uC for 1 h, the cooling rates were measured for AC and SSC conditions using a Flux Hydra data logger, which recorded time and temperature as the specimen cooled. The air cooling rates of the testpiece blanks (15625650 mm) were y1.39 K s21. Slower rates were obtained by cooling in a baffled reflective stainless steel cylinder, 0.56 K s21. Tensile testpieces 3 mm in diameter with a gauge length of 10 mm were machined from the as-received blanks, the air cooled (AC, 1.39 K s21) blanks and the stainless steel cylinder cooled (SSC, 0.56 K s21) blanks. An Instron universal machine (model 1185) was used for mechanical testing, with a crosshead displacement rate of 2 mm min21. Tests were carried out in the range of temperature from ambient to 450uC. The specimens were heated to the testing temperature in a tube furnace controlled by three K type thermocouples. One of these was placed in the top part of the furnace while the other two were placed in the middle and bottom parts, to control the temperature throughout the length of the furnace. A K type thermocouple placed against the centre of the specimen gauge length was used for temperature measurement and control. Temperature variations during testing were found not to exceed ¡2 K. Each specimen was held for 10 min before testing. Load versus displacement information was converted into engineering stress versus strain curves. The following Materials Science and Technology

:

:

:

:

:

:

:

:

:

:

:

: (2)

where: ia is the mean linear intercept grain size of the ferrite, , is the total length of the measurement line, fp is the volume fraction of pearlite and na is the number of ferrite grains cut by the intersecting line.

Results Flow curves for as received, SSC (0.56 K s21) and AC (1.39 K s21) specimens tested between room temperature and 450uC at a crosshead displacement rate of 2 mm min21 are shown in Fig. 2. Serrated flow, one of the characteristics of DSA, was clearly observed at certain temperatures for as received specimens and specimens cooled at the different cooling rates. Serrations first appeared at y200uC, with their magnitude and frequency increasing as the temperature increased to 250uC. Nevertheless, the serrations were observed to disappear beyond 350uC. The AC (1.39 K s21) specimens exhibited serrations at 200, 250 and 300uC. The SSC (0.56 K s21) specimens also showed serrations at 250 and 300uC, while the as received specimens showed serrations at only 250uC. Table 1 gives the hot tensile test results for as received, SSC (0.56 K s21) and AC (1.39 K s21) specimens. It is noted that UTS, work hardening index n and initial work hardening rate showed an increase between 200 and 350uC. The AC (1.39 K 21) specimens also exhibited a further increase in work hardening index and in initial work hardening rate after testing at 450uC. However, LYS of the specimens fell sharply with increasing test temperature in the range studied (ambient to 450uC), while percentage elongation to fracture decreased until 300uC, after which it increased. The effect of temperature on the UTS is presented in Fig. 3. In the temperature range where serrated flow occurred, the UTS increased gradually with increasing temperature and reached a maximum at 300uC under as received, AC and SSC conditions. The UTS then decreased with a further increase in temperature. It can be seen in Fig. 3 that AC (1.39 K s21) and SSC (0.56 K s21) specimens showed bigger increases in UTS compared with as received specimens. Fig. 4 shows the effect of testing temperature on the flow stress at strains of 0.01, 0.02 and 0.05 for vanadium steel under as received, SSC and AC conditions. It is evident that

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Gu¨ndu¨z and Cochrane Dynamic strain aging of V microalloyed steel 3

3 Ultimate tensile stress (UTS) variation with testing temperature

the steel exhibited an increase in flow stress at temperatures considerably above the normal blue brittle range.

Discussion DYNAMIC STRAIN AGING

a as received, location D2; b stainless steel cylinder cooled (SSC) (0.56 K s21), location D1; c air cooled (AC) (1.39 K s21), location A2

2 Load – extension curves for given specimens

Studies of the effect of temperature on the flow stress (0.2% proof stress) of normalised niobium steels indicate a rapid decrease in strength from room temperature to y100uC, then a slower decrease corresponding to dynamic strain aging from carbon which persisits up to about 275 – 300uC. Thereafter, the changes in flow stress are small or negligible.3 – 7,10,11 This behaviour corresponds to the least and greatest amount of free or uncombined nitrogen. Conventional wisdom indicates that dynamic strain aging owing to nitrogen occurs at y100uC, and that at 250uC is caused by carbon. Any change in the binding energy between interstitials and substitutional elements might be detectable by a change in the dynamic strain aging contribution.

Table 1 Tensile test results* for as received specimens from location D2, stainless steel cylinder cooled (SSC) (0.56 K s21) specimens from D1 and air cooled (AC) (1.39 K s21) specimens from A2 Testing temperature, uC As received 25 100 200 250 300 350 400 450 SSC (0.56 K s21) 25 100 200 250 300 350 400 450 AC (1.39 K s21) 25 100 200 250 300 350 400 450

LYS, MPa

UTS, MPa

Elongation to fracture, %

Work hardening index n¡SD

Initial work hardening rate ds/de¡SD, MPa

Grain size{ ¡SD, mm

468 476 423 393 371 320{ 312{ 303{

607 572 552 558 574 535 504 443

34 35 30 30 34 40 43 34

0.131¡0.002 0.129¡0.002 0.130¡0.002 0.154¡0.003 0.188¡0.003 0.197¡0.007 0.178¡0.004 0.152¡0.006

27.3¡3.8 21.4¡2.3 24.4¡2.2 31.4¡3.2 38.3¡3.4 43.6¡3.1 39.1¡2.0 32.2¡3.5

6.4¡0.20 5.4¡0.16 6.2¡0.19 6.1¡0.19 5.5¡0.17 6.1¡0.19 5.9¡0.18 6.0¡0.19

418 379 364 316 324{ 257 250{ 251{

560 512 500 503 533 525 516 463

40 38 38 32 36 45 43 40

0.154¡0.002 0.159¡0.003 0.168¡0.004 0.189¡0.003 0.195¡0.003 0.244¡0.006 0.248¡0.008 0.207¡0.006

23.1¡1.4 22.3¡1.4 22.7¡1.1 32.8¡2.3 34.7¡3.1 40.5¡3.3 47.3¡4.1 37.9¡3.8

6.1¡0.19 5.7¡0.18 6.5¡0.20 6.0¡0.19 5.8¡0.18 6.7¡0.21 5.9¡0.18 6.3¡0.20

429 394 361 349 311 270 305{ 251{

570 518 500 536 553 550 507 492

42 37 31 36 35 44 44 40

0.139¡0.003 0.143¡0.003 0.158¡0.004 0.188¡0.003 0.225¡0.005 0.244¡0.006 0.192¡0.005 0.226¡0.007

22.5¡2.2 22.1¡1.3 22.7¡1.1 30.7¡2.2 39.2¡2.9 42.6¡3.1 35.6¡3.3 42.6¡3.8

4.9¡0.15 5.6¡0.18 6.1¡0.19 5.8¡0.18 5.6¡0.18 6.3¡0.20 6.3¡0.20 6.5¡0.20

*LYS is lower yield stress, UTS is ultimate tensile stress. {Mean linear intercept. {0.2% Proof stress.

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a as received; b SSC; c AC

4 High temperature tensile properties of steel at given strains

It should be possible, therefore, to detect differences in the state of the carbon or nitrogen by analysis of the dynamic strain aging contribution. Such an analysis must make assumptions about the variation in flow or yield stress with temperature in interstitial free steels. Some studies19,20 have shown that it is possible, using a computer, to fit the temperature dependence of flow stress to an equation of the form s~sg zsp zDszs0 (_e=b)T=T0 :

< = >

:

:

:

:

:

:

: (3)

selecting s0, b and T0 the room temperature value should be that used in the standard structure – property relationship for structural steels. It was found in the present work that values of 1000 MPa, 16109 and 2900 K, respectively, resulted in a room temperature value for the Hall – Petch intercept of 70 MPa, which is consistent with the room temperature ky values used. The parameter ky in the sg term represents the slope of the sy – d21/2 relationship, where sy is the lower yield stress, and is believed to be a combination of stress intensification factors needed to propagate dislocations out of boundaries and transfer yielding from one grain to the next. The structure of the boundary regions is therefore regarded as important in controlling the lower yield strength, and the ease with which dislocations can propagate out of the boundary is believed to be dependent on the impurity concentration.21 Various models have been proposed for ky, but most do not consider the temperature dependence above room temperature. The model formulated by Li22 indicates that ky will be dependent on the details of the strength of the interaction between solutes and grain boundary defects. Russell et al.23 concluded that the presence of carbon (or nitrogen) in the grain boundaries is a necessary condition for the existence of a distinct upper yield point in steel. Tjerkstra24 suggested that dislocations penetrate the grain boundaries in steel and implied that there should be a different ky value in the Hall – Petch equation when carbon and nitrogen are concentrated at grain boundaries. Adair et al.25 found that the ky value for vacuum melted iron, water quenched from 700uC, was lower than that when specimens were furnace cooled or when water quenched specimens were reheated for 1 h at 100 K intervals beginning at 100uC followed by furnace cooling. Reheating or slow cooling promotes a substantial amount of grain boundary segregation. As a result, the upper yield stress, lower yield stress and percentage Lu¨ders strain are seen to increase. However, water quenching gives a more rounded yield drop, a smaller lower yield stress and a smoother, shorter Lu¨ders strain region owing to little or no grain boundary segregation. It is clear from the foregoing discussion that the cooling rate (fast or slow cooling), or reheating after fast cooling, affects the segregation of impurities (solute atoms or alloying elements) to the grain boundaries where they influence the ky value. Therefore, for present materials tensile tested at 250 or 450uC after cooling at different rates, the carbon or nitrogen levels at the boundaries could change, hence altering the ky value. Some indication of the extent to which the value of ky is dependent on temperature was attained by examining the experimental data. Two as received specimens obtained from locations E1 and D2 having mean linear intercept grain sizes of 11.2 and 6.2 mm, respectively, were tensile tested at 25, 250 and 450uC. Table 2 gives the actual lower yield strength sact, calculated room temperature lower yield strength scalc on the basis of ky~18.1 MPa and solid solution hardening from manganese and silicon, precipitation

21/2

where sg is a term associated with grain size d (kyd ) and solid solution hardening from manganese and silicon, sp and Ds are precipitation and cluster hardening, respectively, . s0 is the yield stress, e is the strain rate, T is the temperature and b and T0 are material constants. If the temperature dependence of the grain size term can be established, then it is possible to compare the effects of dynamic strain aging both in the as received (as rolled) condition and after heat treatment by correcting the tensile test data for inevitable differences in grain size, which are a result of processing. However, the temperature dependence of sg, sp, Ds and other contributions to strengthening need to be taken into account. Each term is considered in turn. The final term in equation (3) is, of course, temperature dependent, and in Materials Science and Technology

Table 2 Tensile test results* for as received specimens from location E1 and D2 Testing temperature, uC 25 250 450

Code

sact, MPa

scalc, MPa

sp , MPa

d21/2, mm21/2

ws, MPa

E1 D2 E1 D2 E1 D2

474 468 379 393 333 303

321 378 329 383 334 384

153 90 153 90 153 90

9.4 12.5 9.85 12.8 10.1 12.9

321 378 226 303 180 213

*sact is actual lower yield strength, scalc is calculated room temperature lower yield strength, sp is precipitation strengthening, d is grain size, ws~sact2sp.

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?

@

A

B


C

strengthening sp, grain size d21/2 (mm21/2) and the difference between the actual lower yield strength and precipitation strengthening wsLYS(~sact2sp) at 25, 250 and 450uC for these specimens. The similarity of the lower yield strength values calculated for 25uC suggests that the underlying assumptions are broadly correct. It can be noted from Table 2 that the difference in wsLYS is 77 MPa and in grain size is 2.95 mm21/2, which gives ky~26.1 MPa mm1/2, for E1 and D2 specimens tested at 250uC. However the difference in wsLYS is 33 MPa and in grain size is 2.80 mm21/2, which gives ky~11.8 MPa mm1/2, for E1 and D2 after testing at 450uC. The precipitation strengthening was the assumed to be the same at 250 and 450uC as at room temperature, since the specimens came from the same region of the parent section. The difference owing to grain size at 250uC indicates a rather larger value for ky, 26.1 MPa mm1/2, compared with that at room temperature. Conversely, calculation for 450uC suggests a smaller value than that at room temperature, 11.8 MPa mm1/2. Considering the mechanisms operating, another possibility which could explain the variation in flow stress is more reasonable, in that dynamic strain aging invokes a solute drag force owing to dislocations and interstitials. Therefore, it seems more realistic that the observed change is the result of an additional friction stress rather than a material increase in ky. Following the suggestion of Li,22 a temperature dependence given by ln ky versus 1/T was used to model the change in ky with temperature. To show temperature dependence of ky it is necessary to assume that any precipitation hardening contribution to the flow stress sp is comparatively independent of temperature. This assumption is reasonable because of the small restricted possibility of the climb process.26 However, cluster hardening Ds should be temperature dependent, because the binding energy between interstitials and substitutional elements varies with varying temperature. To illustrate the contribution of dynamic strain aging to yield stress, first, values of predicted lower yield stress were calculated according to equation (3) above. The difference between the actual and the estimated lower yield stress Ds indicates the contribution of clusters or solid solution strengthening to the yield stress at a given temperature. Figure 5 shows the Ds variation with temperature for the vanadium microalloyed as received, AC (1.39 K s21) and SSC (0.56 K s21) specimens. For the as received vanadium treated steel, the increment in lower yield strength owing to dynamic strain aging Ds initially increases up to 100uC then declines to a more or less constant value of – 20 MPa at 450uC. In contrast, the vanadium treated steel after air cooling (1.39 K s21) demonstrates an initial sharp rise in Ds in the temperature range 100 – 250uC, declining thereafter but with a pronounced peak at 400uC. The data for the SSC (0.56 K s21) material shows intermediate effects but with two peaks at 200 and 300uC, the curve at higher temperatures being almost identical in shape to that of the as received vanadium steel. The vanadium microalloyed steel revealed different patterns of initial work hardening rate and work hardening index under the as received, AC and SSC conditions (Table 1). Steel specimens showed higher values of initial work hardening rate and work hardening index at testing temperatures of 200 – 350uC. However, the AC (1.39 K s21) and SSC (0.56 K s21) specimens showed higher initial work hardening rate and work hardening index than those of the as received specimens. The AC (1.39 K s21) specimens also exhibited slightly higher values of these parameters compared with the SSC (0.56 K s21) specimens, owing to the higher cooling rate. In addition, the AC (1.39 K s21) specimens showed a further increment in work hardening index and initial work hardening rate at 450uC.

5 Difference between actual lower yield strength and estimated lower yield strength Ds: Hall – Petch constant ky is assumed to be temperature dependent

FP

The AC (1.39 K s21) specimens also showed an increase in flow stress at 400uC for 0.01 and 0.02 strains (Fig. 4). At larger strains (data for 0.05 shown), the AC (1.39 K s21) specimens exhibited a broad peak extending to 300uC, after which flow stress decreased with temperature (Fig. 4), but the decline in flow stress was much less at 450uC compared with the SSC (0.56 K s21) or as received specimens. It is evident that the AC (1.39 K s21) specimens demonstrated an increase in flow stress at temperatures considerably above the normal blue brittle dynamic strain aging range. Such observations are consistent with dynamic strain aging; however, the published data for aluminium killed steels do not show such pronounced changes in ultimate tensile strength as those illustrated in the present paper: see, for example, the work of Glen.10

RELATIONSHIP BETWEEN DYNAMIC STRAIN AGING AND PRECIPITATION (SOLUBILITY) The solubility product of AlN, VN and VC is 4.661025, 1.661024 and 4.261022, respectively, at 900uC.27 It is clear that the solubility of VC is higher than that of AlN (or VN), and therefore AlN should be present at the austenitising temperature. Also, the solubility product of aluminium nitride is lower compared with vanadium carbonitrides. A value of and 0.011 wt-% aluminium was calculated to be in solid solution at 900uC, which is fairly low compared with that for vanadium. Table 3 lists the amounts of free Al, V, N and C in solid solution at 900uC calculated according to the equilibrium solubility product data of Narita.27 The solubility data of Irvine et al.28 give much lower values for soluble aluminium and nitrogen of 0.0066 and 0.0027, respectively. This would suggest that more aluminium present as AlN (82% compared with 69%) would be expected if the Irvine data were applied to vanadium microalloyed steel, but nevertheless implies that some aluminium remains in solution at an austenitising temperature of 900uC. Table 3 Amount of free Al, V, N and C in solid solution at 900uC* AlN, 900uC ks~4.661025

VN, 900uC ks~1.661024

VC, 900uC ks~4.261022

[%Al]

[%N]

[%V]

[%N]

[%V]

[%C]

0.011

0.0042

0.051

0.0032

0.057

0.128

*ks is equilibrium solubility product.


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The higher cooling rate of the AC (1.39 K s21) specimens lowers the transformation temperature and refines the precipitate particle size. There may not be enough time for complete precipitation in the AC (1.39 K s21) specimens, which may contain some of the interstitial carbon or nitrogen atoms in solid solution. Gu¨ndu¨z29 studied the effect of cooling rate on yield behaviour in Ti – V – Al interstitial free steel, and observed that slow cooling rates tended to result in continuous yielding, indicating removal of carbon and nitrogen interstitials, because slow cooling allowed sufficient time for precipitation of carbonitrides. From this it follows that the AC (1.39 K s21) and SSC (0.56 K s21) conditions create differences in precipitate distribution and different amounts of both substititutional and interstitial atoms in solid solution, all of which affect the work hardening behaviour. It is known that the solubility of VC is dramatically higher than that of any other common microalloy precipitate.30 Therefore, during austenitisation at 900uC, vanadium, carbon and/or nitrogen should be in solid solution in austenite. Vanadium, carbon and/or nitrogen could still be in solid solution in ferrite owing to a faster cooling rate which would suppress the precipitation of microalloyed carbonitrides, as proposed by Gladman.18 Therefore, a further increment in work hardening values may have been caused by the presence of vanadium atoms in solid solution at temperatures where they would make dislocation movement more difficult, according to a substitutional strain aging mechanism advanced by Leslie.31 It can be seen from Fig. 3 that after hot tensile testing at 450uC, the AC (1.39 K s21) specimens showed the largest value of UTS. This suggests that vanadium atoms in solid solution were responsible for an increase in dislocation density during testing at 400 – 450uC. Similar conclusions were drawn by Glen10 and Baird,3 – 7 who suggested that such increases are not consistent with substitutional diffusion to dislocations but with substitutional – interstitial interactions. It is reasonable, given the mechanism of VC precipitation in ferrite, that clusters of vanadium and carbon atoms could develop if the ferrite/austenite interface moves too rapidly for precipitates to form, thereby pinning the dislocations. These clusters would, in effect, be ‘frozen in’ during the phase transformation, and thus cause differences in work hardening at high temperatures in the same way as proposed by Baird and Jamieson12 for Fe – Mn – N alloys. They studied the effects of separate and combined additions of 1.6%Mn and up to 0.04%N on the tensile properties of iron in the range 20 – 600uC. At all temperatures, nitrogen additions gave greater increases in yield stress in the presence of manganese than in its absence, and at high temperatures the rate of strain hardening was also greater. These effects are ascribed partly to a solid solution hardening mechanism caused by pairs or small clusters of manganese and nitrogen atoms; this hardening persists up to much higher temperatures than solid solution hardening resulting from nitrogen alone. A second effect is that dynamic strain age hardening owing to nitrogen persists to higher temperatures in the presence of manganese because of the reduction in mobility of nitrogen atmospheres around moving dislocations. It is mentioned above that vanadium microalloyed steel occasionally shows low toughness in the subcritical HAZ (SCHAZ) regions. While no structural failures have been observed to result from this phenomenon, low and variable values at the commonly specified test location, for example FLz3, FLz5 mm, (FL is fusion line), create problems in meeting specifications. This may be attributed to dynamic strain aging, because the present results indicate that the high temperature tensile behaviour of a vanadium microalloyed steel depends on the presence of clusters believed to be of vanadium and carbon atoms. The materials showed Materials Science and Technology

a larger work hardening index and a rapid decrease in strength from room temperature to 100uC, then a slower decrease corresponding to dynamic strain aging from substitutional vanadium and interstitial carbon atoms under conditions relevant to the SCHAZ. Therefore, it can be concluded that the interaction between dislocations and clusters of vanadium and carbon atoms at 200 – 450uC may decrease toughness and the impact transition temperature and lead to subcritical embrittlement in the SCHAZ.

Conclusions Dynamic strain aging has been studied in a vanadium microalloyed steel. The following conclusions emerge from the study. 1. Air cooled (AC) (1.39 K s21) and stainless steel cylinder cooled (SSC) (0.56 K s21) conditions create differences in the amounts of both substititutional and interstitial atoms in solid solution, all of which affect the high temperature tensile behaviour of the steel. 2. Dynamic strain aging takes place in the steel during tensile testing in the temperature range from 200 to 350uC at a crosshead displacement rate of 2 mm min21. As a result, the ultimate tensile stress and initial work hardening rates exhibit maximum values. 3. The AC specimens are more susceptible to dynamic strain aging than SSC or as received specimens as evidenced by increased serration at higher temperatures and by a slightly higher increment in ultimate tensile strength. The vanadium steel also shows an increment in flow stess at 0.01 and 0.02 strains at a higher temperature (400uC) in the air cooled condition. In addition, AC specimens of vanadium steel reveal a higher increment in flow stress at 0.05 strain at temperatures considerably above the blue brittle range. 4. Vanadium in solid solution with carbon and nitrogen, obtained by cooling in air, removes the serration behaviour at higher temperatures. 5. The high temperature tensile behaviour of the vanadium microalloyed steel depends on the presence of clusters believed to be of vanadium and carbon atoms. The interaction between dislocations and clusters of vanadium and carbon atoms at 200 – 450uC contributes to dynamic strain aging, which may decrease the toughness and impact transition temperature and lead to subcritical embrittlement in the SCHAZ.

Acknowledgement The authors would like to thank the University of Leeds, Department of Materials for the provision of research facilities.

References 1. j. d. baird: Metall. Rev., 1971, 16, 1 – 18. 2. j. k. chakravartty, s. l. wadekar, t. k. sinha and m. k. asundi: J. Nucl. Mater., 1983, 119, 51 – 59. 3. j. d. baird: Iron Steel, 1963, 36, 186 – 192. 4. j. d. baird: Iron Steel, 1963, 36, 326 – 334. 5. j. d. baird: Iron Steel, 1963, 36, 368 – 374. 6. j. d. baird: Iron Steel, 1963, 36, 400 – 405. 7. j. d. baird: Iron Steel, 1963, 36, 450 – 457. 8. w. c. leslie and r. l. rickett: J. Met., 1953, 5, 1021 – 1030. 9. p. deb, v. raghavan and v. ramaswamy: Trans. Indian Inst. Met., 1981, 34, 53 – 58. 10. j. glen: J. Iron Steel Inst., 1957, 186, 21 – 48. 11. j. glen: J. Iron Steel Inst., 1958, 195, 114 – 135.

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12. j. d. baird and a. jamieson: J. Iron Steel Inst., 1966, 204, 793 – 803. 13. c. g. interrante: Proc. ‘Dynamic fracture toughness process’ Conf., London, 1976, 327 – 338. 14. j. f. lancaster: ‘Metallurgy of welding’, 5th edn; 1994. 15. r. e. dolby and g. g. saunders: Natl. Semin., 1974, 43 – 66. 16. w. b. morrison, r. c. cochrane and p. s. mitchell: ISIJ Int., 1993, 33, 1095 – 1103. 17. g. e. dieter: ‘Mechanical Metallurgy’, 287; 1988, New York, McGraw – Hill. 18. t. gladman: ‘The physical metallurgy of microalloyed steels’, 1997, London, The Institute of Materials. 19. w. c. leslie: Metall. Trans., 1972, 3, 5 – 26. 20. r. sanstrom and y. bergstrom: Met. Sci., 1984, 18, 177 – 186. 21. b. mintz: Met. Technol., 1984, 11, 52 – 60. 22. j. m. li: Trans. AIME, 1963, 227, 239 – 247.

23. t. l. russell, d. s. wood and d. s. clark: Acta Metall., 1961, 9, 1054 – 1060. 24. h. h. tjerkstra: Acta Metall., 1961, 9, 259 – 263. 25. a. m. adair, r. e. hook and r. l. mcgaughey: Trans. AIME, 1966, 236, 174 – 178. 26. a. kelly and r. b. nicholson: ‘Strengthening methods in crystal’; 1971, London, Elsevier. 27. k. narita: Trans. Iron Steel Inst. Jpn, 1975, 15, 145 – 152. 28. k. j. irvine, f. b. pickering and t. gladman: J. Iron Steel Inst., 1967, 205, 161 – 182. 29. s. gu¨ndu¨z: ‘The effect of cooling rate on the yield behaviour in Ti – V – Al interstitial free steels’, MSc thesis, Zonguldak Karaelmas University, Karabu¨k, Turkey, 1996. 30. t. gladman: Ironmaking Steelmaking, 1989, 16, 241 – 245. 31. r. m. jamieson and r. kennedy: J. Iron Steel Inst., 1966, 204, 1208 – 1210.


May 2003 Vol. 19

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Authors Queries Journal: Materials Science and Technology Paper: 5502 Dear Author During the preparation of your manuscript for publication, the questions listed below have arisen. Please attend to these matters and return this form with your proof. Many thanks for your assistance

Query Reference

Query

1

Author: Is ‘where K is a constant’ OK?

2

. Author: Is e in Equation (3) OK?

3

Author: Is the definition of ‘s0’ OK?

4

. Author: Is ‘e’ OK?

5

Author: Is ‘…in the present work’ OK?

6

Author: Is the definition of ‘sy’ OK?

7

Author: Is the sense of sentence beginning ‘Adair et al. found that...’ OK?

8

Author: Is ‘…for present materials’ OK?

9

Author: Is Table 1 citation OK here?

10

Author: Is it OK to replace ‘vanadium caronitrides’ with ‘V(C,N)’? Or is this repetition of the above?

11

Author: Is ‘wt-%’ OK here?

12

Author: Is ‘vanadium microalloyed steel’ correct?

13

Author: Ref. 31 is Jamieson and Kennedy. Please give correct reference.

14 15

Remarks

Author: Is ‘…present result’ OK? Author: Is point 4 correct? Point 3 above suggests the opposite for AC specimens.


16

Author: In Ref. 13, please give month and sponsor/organiser of the conference.

17

Author: In Ref. 14, please give publisher and publisher location.

18

Author: In Ref. 15, please give seminar title, location, month and sponsor/organiser.

19

Author: Is Ref. 17 correct?

20

Author: In Ref. 19, is it Metall. Trans. A or B?

21

Author: In Ref. 29, is ‘Zonguldak Karaelmas University, Karabu¨k, Turkey’ correct?

22

Author: Is ‘Hall – Petch constant…’ in Fig. 5 caption OK?
