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Journal of Nondestructive Evaluation, Vol. 22, No. 2, June 2003 (© 2003)

Effect of Thermomechanical Treatment on X-Ray Elastic Constants and Residual Stresses in HSLA-100 Steel Aruna Bahadur,1,2 B. Ravi Kumar,1 and S. Ghosh Chowdhury1 Received 1/16/02; Revised 4/21/03

The thermomechanical treatment of HSLA-100 steel is undertaken to upgrade its properties. The plastic deformation in our study was done below and above the recrystallization temperature of austenite, as well in the two-phase (␥ ⫹ ␣) region, followed by accelerated cooling. The plates were subsequently aged at two temperatures. The optical microstructures and hardness were done, and the effect of finish rolling temperature was established. The residual stress, X-ray elastic constants (XEC) and relative peak intensity were determined using X-ray diffraction. The effect of aging temperature on hardness, residual stress, Young’s modulus, and Poisson’s ratio was established. The residual stress values calculated using the XEC values obtained were found to be on the lower side compared to stress values using bulk elastic constants. The maximum difference in stress values is about 27%, which justifies the determination and use of XEC to obtain absolute values of stress.

KEY WORDS: Thermomechanical treatment; HSLA-100; X-ray elastic constants; X-ray diffraction; residual stress; aging temperature; hardness; relative peak intensity.

1. INTRODUCTION

Strength in these steels comes from a highly dislocated aged martensite or bainite microstructure and the precipitation of Cu particles.(3) These steels can be used in as-rolled and aged, normalized and aged, or quenched and aged conditions and can provide several strength and toughness combinations over a wide range of plate thickness.(4) Peak aging temperature for HSLA-100 for maximum yield strength is 454 ⬚C.(5) The best combination of strength and toughness is obtained at 621 ⬚C for this steel.(5) Recrystallization temperature (TR) is found to be 900 ⬚C for this grade of steel.(2)

High-strength, low-alloy (HSLA) steels are modern engineering materials with higher strength and toughness at low temperatures with respect to carbon steels.(1) The main use of the microalloyed HSLA steels has been in the form of hot rolled plate and strip. HSLA-100 steel is basically used for the fabrication of naval ship structures, which are subjected to a complex spectrum of loads and service environments. It is low C (0.06% max) alloy steel that utilizes Cu as a major strengthening element. The microstructure after heat treatment is primarily a martensite-and bainite-like constituent or a combination of both, depending upon the cooling rate/plate thickness and carbon content. (2)

1.1. Thermomechanical Treatment of HSLA Steels

1

National Metallurgical Laboratory, Council of Scientific and Industrial Research, Jamshedpur, India. 2 Corresponding author: E-mail: [email protected]

Two primary objectives of thermomechanical treatment (TMT) are to increase strength by grain 53 0195-9298/03/0600–0053/0 © 2003 Plenum Publishing Corporation

54 refinement and precipitation strengthening. These two depend upon interplay of three parameters: microalloying, plastic deformation, and controlled cooling. All these processes interact to determine the final precipitation-strengthened ferritic microstructure. Retardation of austenite recrystallization is the most effective mechanism for grain refinement. There are four stages of TMT(6), as follows:

1.1.1. Stage I: Deformation Below TR: Controlled Rolling Low-temperature controlled rolling usually involves rolling the steel in a temperature range bounded by T R and Ar 3 , which defines the single phase controlled rolling (CR). Ar3 is the temperature at which ␥ starts to transform to ␣ on cooling. The primary grain refinement in controlled rolling is due to the recrystallization of austenite during hot deformation, leading to dynamic recrystallization. Low-temperature controlled rolling requires high (⬎60%) deformation below TR. The low finishing temperatures cause production losses and frequent occurrence of mixed grain structure and nonuniform properties throughout a single plate.(7) A high TR is a prerequisite for controlled rolling.

1.1.2. Stage II: Deformation Above TR: Recrystallization Controlled Rolling Recrystallization controlled rolling (RCR) is conducted in temperature range of austenite recrystallization. For a given austenite grain boundary area per unit volume (Sv) the ferrite grain size is finer when produced from unrecrystallized austenite than from recrystallized austenite.(8) The steel rolled in the austenite region leads to grain refinement of austenite by static recrystallization after every pass. The temperature difference (TGC – TR) (TGC is the temperature above which austenite grain coarsening takes place) defines the RCR processing.(9,10)

1.1.3. Stage III: Deformation in the Two Phase (␥ ⫹ ␣) Region On rolling in the two-phase austenite and ferrite (␥ ⫹ ␣) region, austenite deforms and partially transforms to ferrite, so that before cooling, there is a

Bahadur, Kumar, and Chowdhury mixture of deformed austenite grain and polygonal ferrite grains.

1.1.4. Stage IV: Controlled Cooling The purpose of Controlled Cooling (CC) is to produce the optimum austenite to ferrite transformation temperature so that the finest ferrite grain size is achieved together with optimum strengthening by microalloy carbide/nitride precipitation. To refine ferrite grain size for a given austenite grain size is to decrease the transformation temperature.

1.2. Residual Stress in HSLA Steels Heat treatments, plastic deformations, or similar texture influencing processes always lead to a change in the residual stress (RS) distribution in the material. In HSLA steels, ferrite is the major constituent; martensite and retained austenite are secondary phases. The transformation of austenite to martensite is accompanied by volume expansion. The change in dimension also depends on the temperature at which the transformation takes place. The layer of martensite nearest to the quenched surface always remains under compressive RS. The remaining phases such as ferrite and retained austenite (RA) are therefore speculated to be subjected to tensile residual stresses to make bulk RS zero. For an externally nonstressed body, sum of RS, each weighted by volume fraction of corresponding phase must be zero. In low-carbon steels, bcc martensite has the same lattice size as the fcc ferrite. The difference in diffraction peaks is that martensite has a broad peak because of the presence of high internal microstresses and the ferrite peak is narrow. Superimposed, the two peaks appear as mainly ferrite peak, with broadened tails near the base.

1.3. Bulk Elastic Constants Neither composition nor structure of low-alloy steels has significant effect on Young’s modulus E, except that austenite has a slightly lower value than other structures of same composition at the same temperature.(11) Poisson’s ratio ␯ is generally 0.3 in steels. It is very sensitive to minor changes of E and shear modulus. Structural changes have little effect on ␯, but it increases from 0.28 to 0.31 as temperature increases from

Effect of Thermomechanical Treatment on X-Ray Elastic Constants and Residual Stresses 20⬚ to 850 ⬚C. Austenite has slightly higher values (0.29–0.34).(11) The modulus of elasticity and Poisson’s ratio for HSLA-100 steel was measured by precision methods and the reported values are E ⫽ 196980 MPa and ␯ ⫽ 0.29.(3)

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samples (numbers 1–12) were austenetized at 1100 ⬚C for 60 min and water quenched (WQ). The details of various TMT given to these samples are listed in Table II. Table I. Chemical Analysis of HSLA-100 Sample Element

Wt. %

1.4 X-Ray Elastic Constants The bulk mechanical elastic constants are averaged over all the orientations and should not be used in X-ray diffraction work. Most materials are not isotropic, and hence elastic constants determination by mechanical means can lead to errors in stress measurement of as much as 80%.(12) Different methods of processing also affect their values.(13) This could be due to grain interactions, texture, or local composition variation. Other source of discrepancy between isotropic and measured elastic constant values are surface relaxation(14) and surface preparation. Long wavelengths and/or deep polishing/grinding effects can give erroneous results,(15) not being representative of the bulk material. The X-ray method is a selective one, because stress is measured only in a specific crystallographic plane, having normal to the grains bisects the incident and diffracted beams. It is necessary to determine experimentally the X-ray elastic constants (XEC) for that plane in order to convert measured strains into stress values. The theoretical values may not agree with experimental ones because of dependence of XEC (Young’s modulus E and Poisson’s ratio ␯) on (i) composition, (ii) second phase components, (iii) grain size, (iv) microstructure, (v) deformation, and (vi) heat treatment. The objective is to study quantitatively the effect of thermomechanical treatment in HSLA-100 steel (Table I) on the X-ray elastic constants. The residual stress is evaluated initially using bulk elastic constants available in the literature and later using the XEC constants generated here. This is done to establish the difference between the values of stress and also to obtain accurate values of residual stress. The effect of TMT on microstructure and hardness are also recorded.

2. EXPERIMENTAL 2.1. Thermomechanical Treatment of HSLA Steels From the 6-mm-thick plate available, samples of 100 mm length and 30 mm width were prepared. All

C Cu Ni Cr Mo Al Nb Mn P S Si N Sn V, Ti, Sb, As

0.05 1.23 1.77 0.61 0.51 0.025 0.037 1.00 0.009 0.001 0.34 0.0102 0.008 ⬍ 0.003

Table II. TMT of HSLA-100 (All Samples Initially Austenetized 1100 ⬚C/60 m/WQ) Set no. I

II

III

IV

Sample no. 1 2 3 4 5 6 7 8 9 10 11 12

Hot rolling (temperature/ deformation/WQ)

Reheat (temperature/ time/WQ)

Nil Nil Nil 850 ⬚C/50%/WQ 850 ⬚C/50%/WQ 850 ⬚C/50%/WQ 953 ⬚C/50%/WQ 953 ⬚C/50%/WQ 953 ⬚C/50%/WQ 950 ⬚C/30% + 900 ⬚C/25%/WQ 950 ⬚C/30% + 900 ⬚C/25%/WQ 950 ⬚C/30% + 900 ⬚C/25%/WQ

Nil 490 ⬚C/60m/WQ 650 ⬚C/60m/WQ Nil 490 ⬚C/60m/WQ 650 ⬚C/60m/WQ Nil 490 ⬚C/60m/WQ 650 ⬚C/60m/WQ Nil 490 ⬚C/60m/WQ 650 ⬚C/60m/WQ

Table III. X-Ray Elastic Constants for ␣-Fe (211) plane of TMT HSLA-100 Plates Sp. no. 1 2 3 4 7 8 9 10 12

E211 (MPa)

␯211

186313 217099 200444 245333 149156 156976 140218 213326 192384

0.331 0.327 0.289 0.357 0.352 0.300 0.352 0.280 0.325

56 2.2. Optical Microstructures Small samples were cut from all sheets, mounted on Bakelite and polished on emery paper and cloth, and etched using nital solution (2% nitric acid in ethanol). Optical photomicrographs were taken at a magnification of 500. 2.3. Measurement of Residual Stress (RS) The residual stress on surface of TMT steel was measured by X- ray diffraction. The interplanar spacing serves as an internal strain gage and the reversible change in the distance between atoms is proportional to the stress acting on the planes. Portable stress analyzer with Cr K␣ radiation, 3-mm collimator size and 10-s exposure was used. The equipment was first calibrated with the help of a standard, stress-free Fe powder. All samples were cleaned with 15% HC1 acid for a fixed time. The shift in the ␣–Fe (211) diffraction peak position that occurs at 2␪ ⫽ 156⬚ was measured accurately using sin2 ⌿ technique. Here ⌿ denotes the angle between normals to the sample surface and the reflecting planes. The sample is oriented at four different ⫾ ⌿ angles to the incident X-ray beam. A series of peak shift measurements were made with planes of atoms oriented at different angles with respect to the sample surface and hence the surface stresses. A straight-line relationship was observed in the plot between peak shift and sin2 ⌿. The slope of the plot m, along with Young’s modulus E ⫽ 196980 MPa and Poisson’s ratio v ⫽ 0.29 for the diffraction plane used, gave the value of the surface stress ␴␸ as follows: m ⫽ ␴␾ (1 ⫹ v)/E 2.4 X-ray Elastic Constants (XEC) Determination The specimens were cut to prescribed sizes of 34mm length, and surfaces were prepared using the standard procedure specified. The strain gages and tabs were pasted just off center on the sample to avoid X-rays falling on them. The four-point bend fixture designed and fabricated at our laboratory has the lower and the higher spans of 20 mm and 88 mm, respectively. The lead wires were soldered and balanced in the quarter bridge configuration in the strain meter. The servohydraulic material testing machine Instron 8502 was calibrated for the 5-ton load cell. The samples (numbers 1–12) were loaded in steps, up to 75% of yield strength, and strain readings were recorded while loading and unloading. To improve the statistics of measurements, a

Bahadur, Kumar, and Chowdhury minimum number of five readings were taken on each sample during loading and unloading cycles and five number of samples were taken for each condition. The applied stresses ␴ were calculated as follows:

␴ ⫽ 3Pa/bh2 where b is breadth, h is thickness, a is arm length and P is load. On the stress analyzer, the samples were fitted in the four-point bend fixture. Loads were applied by screw-driven mechanism to obtain the same strain values. At all selected strain levels, the sample was exposed to Cr K␣ X-rays (30 kV, 6 A) at ⌿ ⫽ 0 and ⫾ 45⬚ at room temperature of 24 ⬚C. From the plots between interplanar spacing d and applied stress, the slopes ⌿ ⫽ 0 and ⫾ 45⬚ (m␺⫽o & m␺⫽45), the values of XEC were obtained as follows: Young’s modulus Ehkl ⫽ (d⬜)/(2m␺⫽45 – m␺⫽0) Poisson’s ratio vhkl ⫽ (⫺m␺⫽0)/(2m␺⫽45 – m␺⫽0) 2.5. Relative Peak Intensity by XRD Using Co K␣ radiation, diffractograms of all samples were prepared from 2␪ ⫽ 45 to 115⬚ at scanning speed of 2⬚/min. The relative intensities of all ferrite peaks were determined and plotted.

3. RESULTS & DISCUSSION 3.1. Microstructures and Hardness in Quenched and Aged HSLA Steels The optical microstructures observed in this study under different TMT conditions show plate-shaped products, which may be bainite/acicular ferrite. Figure 1 is a representative plot. Microstructures of other samples show that plate-type morphologies are breaking down with precipitates along the grain boundaries, as well as within the grains. Figure 2 shows the macro hardness plot of TMT HSLA plates as a function of aging temperature. The results are described in the following section.

3.1.1. Effect of Finish Rolling Temperature In sets II–IV, samples were quenched directly after deformation from finish rolling temperature (FRT). The hardness of samples without aging is found to depend

Effect of Thermomechanical Treatment on X-Ray Elastic Constants and Residual Stresses

Fig. 1. Optical photomicrograph of sample 1 (I set), austenitized 1100 ⬚C/1 h/WQ, aged 490 ⬚C/1 h/WQ, shows polygonal grains and precipitates inside grains and on grain boundaries X500

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In this study, the first set of samples (virgin) was quenched directly from austenetizing temperature. No deformation was imparted to these samples. The hardness in sample 1 is low because it is completely due to solid solution hardening. However, the hardness in sample 7, quenched from above TR, is even less than that of sample 1. This could be due to the precipitation of alloying elements at T R , which is lower than the autenitizing temperature for sample 1, and hence reduce the effect of solution strengthening. The contribution from precipitation hardening in sample 7 is not substantial because the precipitation formed in the austenite temperature range is considered too coarse to contribute much to precipitation hardening. The FRT reportedly (2) did not have a dramatic effect on the hardness response of TMT HSLA-100 plates, as is observed in the present studies.

3.1.2. Effect of Aging Temperature

Fig. 2. Hardness as a function of aging temperature of TMT HSLA-100 plates.

inversely on FRT, as shown in Figure 2. The samples quenched from progressively lower FRT, such as 7, 10, and 4, would lead to increasing precipitation, thereby producing higher hardness upon quenching as a result of precipitation hardening in addition to solution hardening. The precipitation hardening results principally from fine austenite-ferrite interphase precipitation and precipitation in ferrite phase. In this class of steel, solid solution strengthening is much less effective compared to precipitation strengthening.

The hardness in all sets is the lowest in plates without aging condition, such as sample numbers 1, 4, 7, 10, compared to the aged samples (Fig. 2). The trend in set I, (Fig. 2) shows two typical hardness peaks expected for conventional HSLA plates. The first hardness peak is generally observed around 450 ⬚C because of the precipitation of Cu. The second peak observed around 665 ⬚C is attributed to the transformation products of the newly formed austenite because of the low Ac1 temperature, which denotes the lower transformation temperature. In sets II–IV, aging at 650 ⬚C presumably produces further fine Cu precipitate, which increases the hardness in samples 5, 8, and 11 with respect to unaged samples. In sets II–IV, aging at a higher temperature of 650 ⬚C leads to decline of hardness in hot rolled samples 6, 9, and 12 because of growth of precipitates and over aging.(5) Our observations are in line with published work,(2) which shows hardness versus aging temperature for 1/2⬙ thick HSLA-100 steel with interrupted accelerated cooling under different thermomechanical control processing. This general aging response may be considered typical for HSLA-100.(5,16)

3.2. Residual Stress (RS) in HSLA Steel The RS determined in ␣ -Fe with (211) plane (E ⫽ 196980 MPa, v ⫽ 0.29) in this work is plotted in

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Bahadur, Kumar, and Chowdhury of 650 ⬚C relieves the stress practically completely as in samples 3, 6, 9 in Figure 3. Thermal stresses appear to be more significant than hot deformation and transformation stresses.

3.3. Young’s Modulus in HSLA-100 A representative plot between applied stress and interplanar spacing for (211) plane of ␣-Fe is shown in Figure 4 for HSLA plate 1. From the slopes of these straight lines, the values of XEC were evaluated as given in Table II. Young’s modulus E211 was determined from the slopes of above plates for all conditions of TMT; the values are plotted in Figure 5, as follows: E211 varies between 149,000 to 245,000 MPa, which encompasses the reported bulk values of 196,980 MPa for HSLA-100(3) and 211,000 MPa for ␣-Fe. ii. The change in E 211 is significant between various sets, but not within different aging temperatures. iii. E211 is generally correlatable with hardness. The value of E211 increases as hardness increases in sets I, II, III, and IV Figures 2 and 5. iv. E211 is of the same level when no deformation (set I) and rolling below < T R (set II). The lowest values of E211 were found for set III i.

Fig. 3. Residual stress as a function of aging temperature of TMT HSLA-100 plates.

Figure 3. When steel is cooled from a high temperature, the surface cools faster than the interior and it tends to contract, which is resisted by the interior. The surface is thereby placed under tension. This is balanced by compressive stresses in the interior. The stress, as expected, is found to be tensile on the surface, revealed by pickling of TMT HSLA samples.

3.2.1. Effect of Quenching/FRT Temperature The RS is highest when quenched from austenitizing temperature of 1100 ⬚C because of the introduction of high thermal severity resulting from the largest temperature difference (e.g., sample 1, Fig. 3). The samples quenched from progressively lower FRT, such as 7 and 4, predictably show less RS because of progressively less temperature difference.

3.2.2. Effect of Aging Temperature Aging at 490 ⬚C relieves stress partially, as shown in samples 2, 5, and 8. Aging at a higher temperature

Fig. 4. Applied stress versus d211-values for HSLA sample 1 (R is correlation coefficient of linear scale fit by regression analysis).


Fig. 5. Young’s modulus determined for (211) peak of ␣–Fe for TMT HSLA samples.

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Fig. 6. Poisson’s ratio determined for (211) peak of ␣–Fe for TMT HSLA-100 samples.

3.5. Effect of Texture (RCR), which corresponds to the lowest hardness, as in Figure 2. When deformation takes place below < TR or partly below TR, the structure containing more dislocations is retained. On the other hand this phenomenon is less prevalent on RCR above > TR because of dynamic recrystallization. Therefore hardness, as well as E211, in set III should be lower, as observed here. Moreover, during deformation above TR, strain-induced precipitation would take place, which leads to a change in chemistry of the material on quenching compared to the chemistry during other deformation conditions. This would lead to a larger change in E 211 in set III compared to other sets where deformation takes place at lower temperatures.

3.4 Poisson’s Ratio in HSLA-100 v211 values determined for the (211) plane of ␣-Fe are shown in Figure 6: i.

They vary between 0.29 and 0.37, compared to bulk value of 0.29 for HSLA-100. ii. They change with different sets, as well as within a set, for different aging temperatures.

For random orientation of grains, the relative peak intensity for any peak should correspond to the theoretical intensity given in the powder diffraction pattern. The peak intensity changes because of working (e.g., hot rolling). The grains get oriented along certain preferential directions. The peak intensity thereby increases or decreases depending upon grain orientation of a particular plane.(17) This gives a qualitative idea of the presence of texture. From the diffractograms of different samples, the relative peak intensity of (211) plane of ␣-Fe was computed. The values obtained are plotted in Figure 7: i.

Relative peak intensity varies from 10 to 38, compared to theoretical value of 31 using Co K␣ radiation. Most of our values are below theoretical value because of the effect of TMT. ii. The data on v211 is also correlatable with texture in terms of the difference in diffracted relative peak intensity of the (211) plane of ferrite from the random value, for unaged samples such as 1, 4, 7, and 10 and (Fig. 8). Accordingly, when intensity of (211) increases, more plastic deformation is possible and v211 would decrease. iii. The lowest value of relative intensity is for sample 4 (set II), which predictably shows the highest value of v211.

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Bahadur, Kumar, and Chowdhury

Fig. 7. Relative peak intensity of ␣–Fe (211) by XRD for TMT HSLA-100 plates.

Fig. 9. Relation between Young’s modulus and difference in relative peak intensity of ␣–Fe (211) from random value for unaged TMT-HSLA samples.

as 1, 4, and 7 corresponds closely to the E211 plot in the same samples (Fig. 9).

3.6. Revised Residual Stress Using XEC The values of residual stress were recalculated using the values of XECs determined in this work. The stress values calculated with bulk elastic constants (old) and XECs obtained in each case (revised) are plotted in Figure 10: i.

Fig. 8. Relation between Poisson’s ratio and difference in relative peak intensity of ␣–Fe (211) from random value for unaged TMTHSLA samples.

iv. The changes in E211 are due to a change in texture of the (211) plane as a result of TMT. The plot of difference in relative intensity of (211) from random value for unaged samples such

The low stress values obtained (0–38 MPa) in samples (3, 6, 9 and 12) aged at 650 ⬚C, which also happens to be the stress-relieving temperature for steels. ii. The revised stress values (with XEC) are generally on the lower side compared to the old values (with bulk constants). iii. The difference in the old and new values, within a set, increases with the magnitude of stress values. iv. The difference in the residual stress values determined (using bulk and XEC) is plotted in Figure 11 as a function of difference in relative peak intensity from random value. The difference in RS is maximum when the difference in relative peak intensity is maximum


Fig. 10. Residual stress values determined with bulk elastic constants (old) and X-ray elastic constants (revised).

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loading devices were noted. Cr and Co radiation penetrate steel samples to a sufficiently large depth to yield bulk behavior and are not subject to errors because of surface relaxation. Differences would arise in the case of heavy meals such as tungsten, in which X-ray penetration depth is much less. Their (18) (1 ⫹ v)/E (10⫺6 MPa⫺1) values are 4.97 by X-ray bending, 5.68 by X-ray uniaxial tensile loading, and 5.88 by neutron tensile loading for (211) peak. The X-ray value in the literature(19) is 6.35. The bulk value of the elastic constant is 6.22. The difference between bulk and X-ray bending elastic constant is 20.1% for HSLA ferritic steel.(18) In our study on HSLA-100 steel, the difference between bulk and X-ray bending value for sample 1 (austenetized, undeformed) is 9.01% for the (211) peak. This difference varies with the deformation and is maximum for sample 7. The magnitude of disagreement in elastic constants depends on the state of the materials, and a “constant” XEC may not exist. This necessitates the experimental determination of XEC if one wants to measure RS that originates with deformation. Experimental XEC should be generated for a specimen exactly the same (same composition, grain size, heat treatment, deformation history).

4. CONCLUSIONS The following conclusions are drawn from TMT of HSLA-100: i. ii.

Fig. 11. Relation between difference in residual stress values (bulk, XEC) and difference in relative peak intensity from random value

from random value. This shows that the difference in RS values is directly proportional to the texture present in the plate because of the TMT schedules undertaken. The maximum difference of 27% in sample 7, set III, hot rolled above TR. Working with ferritic HSLA steel, Rudnik et al.(18) demonstrated that XEC determined using X-rays and neutrons, agree within experimental error. Also, no difference in X-ray results resulting from difference in

iii. iv.

v.

The microstructures are composed primarily of baintic/acicular ferrite constituents. The hardness of TMT plates, as expected, increases with decreasing FRT. The lowest hardness of the highest FRT sample is higher than that of virgin sample 1, which was quenched from solution treatment temperature of 1100 ⬚C. The hardness of TMT plates is the lowest in the unaged condition. Samples aged at 490⬚C show precipitation hardening in all plates. Subsequent increase in aging temperature to 650 ⬚C leads to a decrease in hardness. In virgin sample, hardness increases further with aging temperature of 650 ⬚C. Residual stress on the exposed surface of TMT plates is tensile in nature in ferrite phase. It is highest when quenched from the austenitizing temperature of 1100 ⬚C and decreases with lowering of FRT.

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Bahadur, Kumar, and Chowdhury vi.

Residual stress is relieved partly on aging at 490 ⬚C and practically completely on aging at 650 ⬚C vii. Young’s modulus E 211 varies between 149,000 to 245,000 MPa. The values are on the lower side on RCR (> TR), which is correlatable with hardness. It is directly proportional to the texture present in-hot rolled HSLA samples. viii. Poisson’s ratio v211 varies between 0.28 and 0.37. It changes within different sets, as well as within a set, for different aging temperatures. It is directly proportional to the texture present in samples on account of hot rolling. ix. The RS values calculated using XECs are generally on the lower side compared to stress using bulk elastic constants. x. The difference in RS values obtained using bulk and XECs is directly proportional to the amount of texture present in samples because of hot rolling. xi. The maximum difference in the stress values is about 27% in sample 7, hot rolled above TR, justifying the determination and use of XECs to obtain absolute values of stress. ACKNOWLEDGMENTS We wish to record our thanks to Mr. B. A. Lakra and Dr. S. Sivaprasad for help during experimentation. Our thanks are due to Dr. D. K. Bhattacharya, Head, Material Characterization and Professor P. Ramachandra Rao, former Director, National Metallurgical Laboratory for their support and encouragement. REFERENCES 1. B. L. Bramfitt and J. G. Spear, A perspective on the morphology of bainite, Metall. Mater. Trans. A, 21A, p. 817 (1990).

2. E. M. Focht, “Proc. Mater. Solution ‘97 on Accelerated Cooling/Direct Quenching of Steels,” ASM, OH, pp. 23–32 (1997). 3. E. J. Czyryca, Internal Report, DTRC-SME-90/21, David Taylor Research Centre, Bethesda, (1990). 4. M. R. Krishnadev, A. Laasraoui, K. Romhanyi, ‘Processing, Microstructures and Properties of HSLA Steels,’ edited by A. J. DeArdo, p. 261, 1988, TMS, PA. 5. E. G. Hamburg, A. D. Wilson, ‘HSLA Steels: Processing, Properties & Applications,’ edited by G. Tither, Z. Shouhua, p. 241 1992, The Minerals, Metals and Materials Society. 6. T. Tanaka, TMS-AIME Conf. Proc. on HSLA, edited by D. P. Dume and T. Chandra, pp. 6–16 1984. 7. M. E. Natishan, “Mechanisms of Strength and Toughness in a Microalloyed, Precipitation Hardened Steel,” David Taylor Research Centre Rep., Bethesda, SME-88–81 (April 1989). 8. R. K. Amin and F. B. Pickering, ‘Thermo mechanical Processing of Microalloyed Austenite,’ edited by A. J. DeArdo, G. A. Ratz, and P. J. Wray, p. 377, 1982, AIME, Warrendale, PA. 9. Y. Zheng, A. J. DeArdo, R. M. Fix, and G. Fitzsimons, Intl. Conf. “Technology & Applications of HSLA Steels,” ASM pp. 85–94 (1983). 10. S. Zajac, T. Siwecki, B. Hutchinson, and M. Attlegard, Recrystallization controlled rolling and accelerated cooling for high strength and toughness in V-Ti-N steel, Metall. Trans. 22A, pp. 2681–2694 (1991). 11. A. J. Fletcher, Thermal Stress & Stress Generation in Heat Treatment, p. 42, 1989, Elsevier London. 12. P. S. Prevey, A method of determining the elastic properties of alloys in selected crystallographic directions for X-ray diffraction residual stress, Adv. X-ray Anal. 20, pp. 345 (1977). 13. R. M. Zhong, I. C. Noyan, J. B. Cohen, Adv. X-Ray Anal. 29, p. 17 (1986) 14. T. Hanabusa, K. Nishioka and H. Fujiwara, Criterion for the triaxial X-ray residual stress analysis, Z. Metallkde. 74, p. 307 (1983). 15. A. D. Krawitz, The use of X-ray stress analysis for WC-base cermets, Matl. Sci. Eng. 75, p. 29 (1985). 16. A. D. Wilson, E. G. Hamburg, D. J. Colvin, S. W. Thompson, and G. Krauss, “Properties and Microstructures of Copper Precipitation Aged Plate Steels” Proc. Microalloy 88, World Mater. Congr. ASM, pp. 259–275 Sept. 1988. 17. H. J. Bunge and C. Esling, editors, Quantitative Texture Analysis, 1982, Deutsche Gesellschaft fur Metallkunde, Adenauerallee 21,6370 Oberursel 1, Germany. 18. P. J. Rudnik, A. D. Krawitz, D. G. Reichel, J. B. Cohen, Adv. X-Ray Anal. 31, p. 245 (1988). 19. Society of Automotive Engineers, Residual Stress Measurements by X-ray Diffraction; SAE Handbook J784a, 2nd ed, 1971, SAE, New York.