Commission of the European Communities

technical steel research

Properties and service performance

HIC-resistant steel - Structure and composite effects

Commission of the European Communities

echnical steel research

Properties and service performance

HIC-resistant steel - Structure and composite effects

R. E Dewsnap British Steel pic

9, Albert Embankment London SE1 7SN United Kingdom

Contract No 7210-KE/813 (1.7.1986-30.6.1988)

Final report

Directorate-General Science, Research and Development

1990 Otf\S ^ M < ^ >

PARL EUROP. Bibiioih.

N.C./tfCo00S&S<

CL EUR 12959 EN

Published by the COMMISSION OF THE EUROPEAN COMMUNITIES

Directorate-General Telecommunications, Information Industries and Innovation

L-2920 Luxembourg

LEGAL NOTICE Neither the Commission of the European Communities nor any person acting on behalf of the Commission is responsible for the use which might be made of

the following information

Cataloguing data can be found at the end of this publication

Luxembourg: Office for Official Publications of the European Communities, 1990

ISBN 92-826-1903-6 Catalogue number: CD-NA-12959-EN-C

ECSC-EEC-EAEC, Brussels Luxembourg, 1990

Printed in Luxembourg

HIC RESISTANT STEELS - STRUCTURE AND COMPOSITION EFFECTS

British Steel pic

ECSC Agreement No. 7210.KE/813

SUMMARY

A laboratory composite mould design has been developed consisting of sand and heavy metal chill plates for the casting of slab ingots with central segregation similar to that in continuously cast slab products. This technique has been used as the basis for a programme involving a study of the HIC resistance of a range of vacuum melted low sulphur calcium treated PRS (Pearlite Reduced Steels) and CMnMoNb and CMnTiB AF (Acicular Ferrite)/bainite steels. All steels including a conventional 0.1% carbon ferrite-pearlite linepipe composition used as control were controlled rolled to 15 mm thick plate and a full standard mechanical properties and HIC assessment was conducted including single sided HIC testing with continuous hydrogen permeation and ultrasonic C scan monitoring followed by backup metallography and microanalysis studies of HIC sensitive and segregated regions. A limited study has also been conducted to determine the effect of phosphorus content and laboratory simulated submerged arc welding on HIC response.

Tensile properties ranged from Grade X60/X65 in the pearlitic microstructural types to Grade X70/X80 with continuous yielding behaviour and a high strain hardening index in the acicular ferrite steels together with excellent notch toughness. Pronounced centreline HIC was found in the Grade X65 ferrite-pearlite control steel along hard martensite/bainite bands containing manganese contents up to 2.2% with a segregation ratio of 1.8. In contrast the higher strength ultra low carbon high manganese AF/bainite steels and particularly the PRS compositions showed intermediate to high levels of HIC resistance with limited segregation band HIC as a consequence of the lower centreline manganese segregation ratios which were typically less than 1.2. These findings support the claim that a reduction in carbon content has a pronounced beneficial influence in reducing the segregation of solute elements. Typical mean CLR values for the ferrite-pearlite control steel, PRS and AF/bainite steels were 17%, 3% and 4-13% respectively reflecting the generally fewer crack systems and reduced cracking intensity of the low carbon materials. However, surface blistering was more pronounced in the AF/bainite steels and particularly in the Grade X80 C1.9MnTiB alloy. Except for the ferrite-pearlite control steel which showed limited cracking in a 1000 h single sided HIC test in NACE solution with a hydrogen threshold less than 1 ml/100 mg, all other steel types remained crack free. Hydrogen permeation studies on single sided HIC testpieces showed classical behaviour on all except the ferrite-pearlite control steel with a rapid rise in permeation current to a peak value within 100-200 h of start of test followed by a steady decay throughout the remaining test period. Hydrogen breakthrough also occurred in shorter times in the low carbon PRS and AF/bainite steels compared with the higher carbon ferrite-pearlite steel.

A reduction in phosphorus content from about 0.011% to 0.003-0.005% had a pronounced beneficial effect on the HIC resistance of the ferrite-pearlite control steel and the 1.9% manganese CMnTiB AF/bainite steel, reducing mean CLR values from approximately 17% and 13% respectively to less than 3%, comparable with results for the ~0.011% phosphorus PRS steel. ' This is attributable to a reduction in segregation band hardness. Severe surface blistering in the 1.9% manganese CMnTiB steel was also greatly attenuated with a reduction in phosphorus content. The mean microhardness of cracked and uncracked segregation bands in steels containing approximately 0.011% phosphorus was well below the generally accepted threshold level of 300 HV for substantial HIC and typically of the order of 280 HV for the ferrite-pearlite steel and 230-240 HV for PRS and AF/bainite steels. The low phosphorus variants of the ferrite-pearlite and 1.9% manganese CMnTiB AF/bainite steel showed a mean decrease in segregation band hardness of approximately 60 H V compared with the intermediate and high phosphorus equivalents. All steels exhibited good weldability with excellent weld metal toughness and full HIC resistance in HAZ and weld metal.

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CONTENTS

Page

SUMMARY III

LIST OF TABLES VII

LIST OF APPENDICES VII

LIST OF FIGURES IX

1. INTRODUCTION 1

2. EXPERIMENTAL TECHNIQUES 2

2.1 Material preparation 2 2.2 Mechanical properties 3 2.3 HIC testing 3

3. RESULTS AND DISCUSSION 4

3.1 Ingot mould optimisation trials. 4 3.2 Ferrite-Pearlite steel 5 3.3 Low carbon perarlite reduced steel 7 3.4 Acicular ferrite/bainite steels 9 3.5 Effect of Phosphorus 11 3.6 Single sided HIC tests 12 3.7 Weldability studies 14 3.8 Microanalysis and microhardness 15

4. CONCLUSIONS 17

5. REFERENCES 17

TABLES 20

FIGURES 49

APPENDIX 158

LIST OF TABLES

1 Chemical Analysis of Experimental Ingots

2. Nominal Controlled Rolling Schedule for All Steels

3. Ingot Casting Variables

4. Mechanical Properties of Controlled Rolled CMnNb Ferrite-Pearlite Steel - Plate E

5. HIC Test Results - Ferrite-Pearlite Steel - Plate E

6. Comparison of HIC Parameters - Ferrite-Pearlite Steel - Plate E

7. Chemical Composition of Pearlite Reduced Steels

8. Mechanical Properties of Pearlite Reduced Steels

9. HIC Test Results - Pearlite Reduced Steels

10. Chemical Composition of AF/Bainite Steels

11. Mechanical Properties of AF/Bainite Steels

12. HIC Test Results - AF/Bainite Steels

13. Chemical Composition of Phosphorus Variant Steels

14. Mechanical Properties of Phosphorus Variant Steels

15. HIC Test Results - Ferrite-Pearlite Steel Plate M 0.005% P

16. HIC Test Results - Acicular Ferrite Steel Plate N 0.003% P

17. HIC Test Results - Acicular Ferrite Steel Plate O 0.019% P

18. Analysis of Crack Data - All Steels

19. Hydrogen Diffusion Coefficients and Critical Hydrogen Concentrations for Plate E Compared With Commercial Linepipe Steel

20. Welding Procedure for Triple Wire Submerged Arc Welds on 15 mm Plate

21. Notch Toughness Properties of Weld Metal

22. HIC Test Results - Pearlitic Steel Welds

23. HIC Test Results-AF Steel Welds

LIST OF APPENDICES

1. Method of Calculating Critical Hydrogen Concentration for Cracking

VII

LIST OF FIGURES

1. Ingot Mould Designs

2. Positioning of Thermocouples in Chill Plate - Ingot A

3. Test Piece Sampling Procedure

4. Thermal History in Chill Plate and Advance of Solidification Front During Casting of Ingot A

5. Photomacrographs of Ingot Sections

6. Effect of Chill Plate Thickness on Columnar Crystal Growth

7. Manganese Concentration in Central Segregation Zone of Ingots A and D

8. Typical Microstructures and Full Plate Thickness Macrostructure Ferrite-Pearlite Steel - Plate E

9. Typical Non-Metallic Inclusion Distributions Ferrite-Pearlite Steel - Plate E

10. Standard Charpy V-Notch Transition Curves for Controlled Rolled Ferrite-Pearlite Steel -Plate E

11. C Scan Traces Showing HIC in Ferrite-Pearlite Steel - Plate E - BP and NACE Solution

12. HIC Along Pearlite and Bainite/Martensite Bands in Ferrite-Pearlite Steel - Plate E

13. HIC Defining Variable Segregation Band Width Ferrite-Pearlite Steel - Plate E

14. Manganese Profiles Across Cracked Pearlite and Bainite/Martensite Bands Ferrite-Pearlite Steel - Plate E

15. Isometric View Showing Manganese Profile in Cracked Bainite/Martensite Band Ferrite-Pearlite Steel - Plate E

16. Standard Charpy V-Notch Transition Curves for PRS Plate

17. Elastic-Plastic Region of Load-Extension Curve - PRS Plate H

18. Typical Inclusion Distribution in Vacuum Melted Pearlite Reduced Steels - Plate F

19. Typical Microstructures of PRS Plates

20. C Scan Traces Showing HIC - PRS Plates

21. Surface Blisters on PRS Plates

22. Full Cross Section Photographs of Test Piece Sections Showing HIC in PRS Plate F

23. Microstructural Aspects of HIC in PRS Plates F and L

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24. Standard Charpy V-Notch Transition Curves for AF/Bainite Steels

25. Elastic-Plastic Region of Load-Extension Curve - AF/Bainite Steels

26. Typical Microstructures of AF/Bainite Steels

27. Typical Crack Distributions in HIC Testpiece Sections - AF/Bainite Steels

28. C Scan Traces Showing HIC - CMnTiB and CMnMoNb - AF/Bainite Steels

29. Surface Blisters on AF/Bainite Steel Plates

30. Typical Examples of HIC in AF/Bainite Steels

31. Standard Charpy V-Notch Transition Curves Ferrite-Pearlite Steel - Low and High Phosphorus

32. Elastic Plastic Region of Load-Extension Curve Ferrite-Pearlite Steel - Low Phosphorus and AF Bainite Steels Low and High Phosphorus

33. Test Piece Sampling Procedure for High and Low Phosphorus Steels

34. C Scan Traces Showing HIC in Low Phosphorus Ferrite-Pearlite Steel and Low and High Phosphorus AF Steels

35. Typical Surface Blistering in Phosphorus Variant Steels

36. Effect of Phosphorus on Mean CLR of Pearlitic + AF/Bainite Steels

37. Multi Testpiece Single Sided HIC Test Cell

38. Schematic Diagram of Test Loop for Single Sided Cell

39. Single Sided Cell in Position on Ultrasonic Testing Jig

40. Change of pH and H S Concentration During 1000 h Single Sided HIC Test

41. C Scan Traces Showing Crack Development in Commercial HIC Susceptible Steel During Single Sided Test

'42. Crack Area Ratio for Commercial HIC Susceptible Steel Subjected to 1000 h Single Sided Test

43. C Scan Trace Showing Crack Development in Plate E Single Sided Test - Exposed for 936 h

44. Hydrogen Permeation Current for Plate E During Single Sided Exposure

45. Hydrogen Concentration at Centre of Single Sided Test Pieces

46. HIC in Single Sided Test Piece - Plate E

47. Hydrogen Permeation Curve for Single Sided Exposure HIC est

48. Arrangement of Plates for Submerged Arc Welding

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49. Charpy V-Notch Transition Curves for Submerged Arc Weld Metal

50. Typical Surface Appearance of Weld HIC Test Pieces

51. Manganese Profiles in Region of Branched Crack System - Plate E, NACE Single Sided Test Piece

52. Manganese Contour Map in Uncracked Segregate Region - Plate E, NACE Single Sided Test Piece

53. Concentration Profiles Across Centreline Crack

54. Manganese Profiles Across HIC and Uncracked Region of Segregate Band - Plate F Manganese Concentration Map in Region of HIC - Plate L

55. Manganese Contour Map in Branched Region of HIC - Plate G, Test Piece 1C, Section A

56. Manganese Profiles Across Stepped HIC - Plate G, Test Piece 6D, Section C

57. Manganese Concentration Maps in Region of HIC - Plates J and K

58. Manganese Profile at Plate Mid-Thickness Position Plate J - Testpiece 6P69K - Section 1A Plate K - Testpiece 6P70K - Section IB

59. Concentration Profiles Across Centreline and Off Centre Cracks

60. Relationship Between Mean Segregation Band Hardness and CLR

XI

HIC RESISTANT STEELS - STRUCTURE AND COMPOSITION EFFECTS

British Steel pic

ECSC Agreement No. 7210.KE/813

FINAL TECHNICAL REPORT

1. INTRODUCTION

Early and subsequent studies of hydrogen induced cracking (HIC) behaviour in linepipe steels established the importance of non-metallic inclusions and particularly inclusion morphology as a key factor in crack initiation1-^. Consequently, most of the early developments were concentrated on improving steel cleanness in general and reducing the fraction of elongated non-metallic inclusions in particular since these were identified as the principal crack nucleators, becoming effectively hydrogen pressurised notches under sour environmental exposure conditions7. The importance of steelmaking practice and sulphur content was therefore recognised4.7.9. While toughness and tensile ductility properties benefited with the introduction of fully killed steels as replacement for the relatively dirty, balanced or semi-killed steels, it was subsequently established that the globular Type I sulphides in the latter were less damaging to HIC performance under the prevailing BP solution (pH ~5) test conditions than the killed steel Type II lenticular sulphides^. Despite the application of clean steel technology practice, the introduction of controlled temperature rolling aggravated the situation, resulting in a greater degree of inclusion elongation compared with conventional hot rolling*. Thus began an era of steelmaking developments during which tonnage vacuum treatment and desulphurisation techniques resulted in cleanness levels frequently considerably better than hitherto obtainable by laboratory and special secondary refining methods. Typically, sulphur and oxygen contents less than .20 ppm together with inclusion modification treatment, currently based on calcium additions10*12 giving a low volume fraction of morphologically acceptable inclusions is now considered necessary for maximum HIC resistance.

While steel cleanness is a major factor and a prerequisite for improved HIC resistance this did not provide the universal solution. Alloy enrichment problems manifest as A and V segregation in ingot casting and centreline segregation in continuously cast slab, resulted in hard band structures of martensite and bainite13"16 which, together with pearlite bands14-17, are recognised as potent crack propagators. Indeed, some authorities are of the opinion that HIC may be self-nucleating in hard microstructures due to embrittlement effects. Consequently, even in clean steels under mild sour test conditions HIC resistance may be seriously impaired. Although reported to the contrary1^ it is now recognised that all the solute elements segregate to varying degrees and while manganese, and particularly phosphorus, which is reported to be the most potent hardenability promoter after boron, are considered to be the most detrimental, the hardenability effect of the minor element additions may be considerable. Secondary to steel cleanness HIC performance is in general therefore directly related to carbon and alloy content16 effectively curtailing development of the traditional controlled rolled ferrite-pearlite type steels above about Grade X65. Indeed the tendency has been towards lean steel practice with lower carbon and a reduction in manganese content to below 1% coupled with balanced microalloying additions and modified controlled rolling practice to compensate for loss of strength. Nevertheless, strong segregation may still be encountered dependent on casting conditions contrary to the claim that manganese does not segregate below bulk contents of about 1%. It is generally agreed that segregation ratio peaks of up to 1.5-2 for manganese and 10 for phosphorus can be obtained in a centre band of typically 10 to 100 urn wide in plate1^.19. While this has undoubtedly influenced the drafting of specifications accepting relatively long planar HIC there is a growing awareness that energy fields are likely to become increasingly sour and consequently material specifications must eventually reflect these conditions.

Many studies of centreline segregation have been carried out aimed at an understanding and ultimately control of the problem at source20*22. Nevertheless despite improvements in casting practice and new

generation continuous casting machine design there appears to be little prospect of a complete solution to the segregation problem within the foreseeable future. Consequently, a wide range of metallurgical options have been investigated aimed at minimising or eliminating centreline HIC in continuously cast derived products. These are either precasting measures intended to provide a slab having acceptable macrosegregation, achieved entirely by dependence on steel chemistry, or post casting measures in which corrective treatments are applied at the slab or plate stages. Typical of the corrective treatments are plate tempering23f normalising and Q & T4.!*.24, high temperature homogenisation both with and without hot deformation24.25 and direct Q & T18 or interrupted cooling treatments from rolling heat26. Except for interrupted cooling to an intermediate temperature involving no further treatment the majority of the post slab casting corrective treatments have found little favour either for technical or economic reasons.

Since segregation band hardness is dependent on carbon content which partitions strongly to the high alloy segregation zone during transformation attention is being focussed on progressively lower carbon content steels. The principal variants are pearlite reduced steels (PRS) with insufficient carbon for continuous pearlite band formation and ultra-low carbon enhanced alloy content steels having acicular ferrite (AF) and bainitic microstructures with little free carbide. PRS type materials are currently being supplied to severe sour service specifications and when based on controlled manganese content appear to give good HIC resistance in laboratory HIC tests compared with the higher carbon variant ferrite-pearlite controlled rolled steels. The ultra-low carbon AF and bainitic steels appear to be largely experimental although published data suggests substantially better HIC resistance than the ferrite-pearlite steels giving a greater tolerance for high alloy concentration in the segregation' band16.2?. Typically at a segregation band manganese content of 3% with bulk carbon of 0.02% a band hardness less than the threshold level for substantial HIC of 300 HV is claimed giving acceptable resistance in a NACE solution test. The various material and process factors affecting HIC resistance in linepipe steels have been comprehensively reviewed28.

The principal objective of this project was to assess the HIC resistance of a range of controlled rolled ultra-low carbon enhanced alloy content AF/bainitic steels compared with traditional ferrite-pearlite materials. This report describes the various stages of the project including the development of a suitable laboratory slab casting technique giving reproducible centreline segregation, the processing of a number of alloys to strength levels within the range X65-X80, HIC assessment and hydrogen permeation studies, and examination of segregation profiles in relation to compositional variables and HIC development. Finally, the results of a limited HIC assessment of laboratory simulated linepipe seam welds on the ultra-low carbon AF/bainite and ferrite-pearlite steels is described.

2. EXPERIMENTAL TECHNIQUES

2.1 Material Preparation

In view of the sensitivity of HIC to microstructural heterogenieties as discussed previously it was realised that the success of this project clearly depended on the provision of experimental materials in a commercially realistic condition for assessment. Due to the prohibitive cost of experimental production heats since the proposed materials were generally not available commercially the project was based entirely on laboratory melted materials. Therefore in order to generate meaningful data on a broad range of compositional and microstructural variables it was necessary to develop a laboratory technique capable of 'casting in' segregation profiles similar to commercial continuous casting practice with the aim of achieving full centreline macrosegregation. The suitability of the various processing options would then be assessed using standard metallographic techniques and where necessary by HIC and microanalysis studies prior to commencement of the main programme.

2.1.1 Ingot Casting

Two approaches to the problem of achieving centreline segregation were pursued as described in published data29*30. The favoured approach was based on a technique involving intense single sided cooling of a slab ingot giving a controllable solidification front with well defined columnar grains and a

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narrow segregation zone along the length of the ingot. An alternative approach is to base the mould design on a 'dumbell' ingot with a heated or insulated central region and large feeder head and lower shrinkage zone aimed at encouraging intense V segregation. Both principles were considered and the various mould designs employed are illustrated in Fig. 1.

It was initially intended that optimisation of casting conditions would be accomplished using these composite moulds constructed of sand, refractory insulating board and mild steel chill plates and analysing the thermal conditions and resultant ingot macrostructures. These data would then be used to design a water cooled copper mould with which to cast a series of vacuum melt ingots for the main experimental programme. Unfortunately, problems were encountered in providing an additional water cooling supply within the vacuum furnace chamber and consequently the final mould design was based on the use of solid metal chill plates.

For the ingot mould optimisation studies a series of CMn top poured 50 kg air melt ingot were cast with typical composition shown in Table 1, and the macrostructure of each was examined before proceeding to the next stage. Thermal conditions during solidification of the first ingot (Ingot A) were calculated using data from thermocouples embedded in one of the mild steel mould plates and positioned as shown in Fig. 2. This information was used in further ingot casting studies at various degrees of superheat culminating in the optimised mould design E, Fig. 1. This mould configuration was employed throughout the experimental programme.

2.1.2 Plate Rolling

Since little data was available on the processing of ultra-low carbon AF steels all materials were controlled temperature rolled to a basically standard schedule designed to develop X65 strength properties in a 0.1% C, 1.25% MnNb conventional linepipe type composition. This involved a double hold thirteen pass drafting schedule to 15 mm plate from a reheating temperature of 1150C, Table 2. Because of the restricted ingot thickness this necessitated a low finishing deformation which was compensated by reducing the intermediate hold temperature compared with a commercial schedule, giving a 3:1 reduction below 880C with a finishing temperature of 720C.

2.2 Mechanical Propert ies

Tensile and standard Charpy V notch energy and fracture appearance transition curves were determined for all materials on transverse test pieces. The tensile tests were conducted at an extension rate of 5 mm/min on full plate thickn.ess test pieces having a gauge length and width of 50 mm and 38 mm respectively.

2.3 HIC Testing

General categorisation of materials and relative HIC response was determined using the standard BP/Cotton HIC immersion tests n accordance with the procedure described in NACE standard TM-02-84(3D. In addition, single sided exposure testing with integral ultrasonic inspection and hydrogen permeation determination was carried out on selected materials. For the ferrite-pearlite steel and the first AF/bainite steel standard HIC tests were initially conducted in both H2S saturated BP solution (substitute sea water formulated according to ASTM Designation D1141-75 - Reapproved 1980) and NACE solution (5% sodium chloride + 0.5% acetic acid formulated according to TM-01-77). It became evident, however, that the lower carbon variants of the ferrite-pearlite compositions and the AF/bainite materials were highly resistant to HIC in BP solution and subsequent tests on these types of materials were conducted only in NACE solution.

In order to assess the consistency of material from the first vacuum melt and also of material from selected subsequent melts HIC test pieces were prepared from two positions along the plate length as shown in Fig. 3, each alternate test piece being tested in the same solution. Test piece preparation involved a milling and wet surface grinding operation to remove approximately 0.75 mm from each of the four principal surfaces and the final 0.25 mm was removed at the rate of 0.04 mm per pass, the test piece ends remaining in the sawn condition. This was followed by dry grinding on a rotary belt linisher using

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progressively finer grit with a final 320 grit finish. The final test piece sizes were approximately 100 mm long x 20 mm wide x 13.5 mm thick. Immediately prior to immersion test pieces were decreased in warm Genklene, a proprietary chlorinated hydrocarbon which has been found to be more effective than the recommended acetone and the effectiveness of this operation was assessed using the water droplet test.

Testing was carried out at 25C 1.5CC in the appropriate solution for a period of 96 h using a solution volume to test piece surface area ratio of 4.5 ml/cm2. This was followed by hydrogen collection over glycerine for 72 h at a temperature of 45C. Surface blisters were highlighted prior to counting by gentle abrasion of the test piece surfaces using 600 grade emery. Prior to sectioning of the test pieces for crack assessment according to standard procedure, i.e. at the -J-, | and positions from one end, the general level of cracking was determined using an ultrasonic C scan technique. Sections were then polished automatically in the unmounted condition to a 1 pm diamond finish and cracks were measured at a magnification of X30 using a separation value of 0.5 mm to define an isolated crack or crack system. All cracks were included in the assessment except those lying entirely within 1 mm of the test piece upper and lower surfaces and the data are presented as a series of ratios defining crack length (CLR %), crack thickness (CTR %) and crack area described as crack sensitivity ratio (CSR %) as defined in NACE standard TM-02-84.

3. RESULTS AND DISCUSSION

3.1 Ingot Mould Optimisation Trials

3.1.1 Casting

Since intense cooling along the length of the ingot, as described in published data2 9 , was the most favoured approach to achieve intense centreline segregation, the mould design for casting of the first air melt slab ingot was based on solid chill plates along the wide faces with insulating board sides. In view of the limited vacuum melt weight available for casting of the experimental materials the width to thickness ratio of the slab ingots was fixed at 2:1 throughout giving a mean thickness of approximately 100 mm.

The principal variable investigated was chill plate thickness as shown in Table 3 although the degree of superheat was ultimately increased in the light of findings throughout the optimisation trials. Thermal history at various positions in the chill plate during casting of the first ingot is given in Fig. 4 together with computed thermal/metallurgical conditions in the ingot. The break in the temperature curves for Thermocouples 3 and 5 is unlikely to be related to thermal conditions in the ingot and is probably a result of temporary loss of contact between thermocouple and mould plate.

Thermal history in the ingot was computed using a heat transfer model developed at Swinden Laboratories, from which the relationship between solidified shell thickness and time was derived. This indicated a considerably lower rate of growth by a factor of approximately four compared with typical continuous slab casting conditions32. In an attempt to increase this rate and thus maximise columnar crystal growth the mould plate thicknesses were progressively increased for the casting of further ingots.

The single attempt to develop V segregates using a dumbell ingot form30 was unsuccessful and consequently this approach was not pursued.

3.1.2 Macrostructure of Trial Ingots

The macrostructure of each air melt ingot was examined to determine the degree of segregation before proceeding to the subsequent stage. Sectioning of each ingot was carried out along the axis at the mid-width position followed by wet surface grinding and etching in a 20% aqueous solution of ammonium persulphate. From the sequence of macrostructures, Fig. 5, it is evident that the most successful mould configuration to evolve incorporated heavy gauge chill plates along the wide faces with sand insulated ends and base and a relatively large feeder head. This latter aspect was a compromise between adequate liquid metal feeding and maximising the volume of useful metal in the ingot.

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Ingot A which was cast using 30 mm thick chill plates showed relatively short columnar crystal growth with a broad central equiaxed zone containing dispersed segregates resembling semi-macro or spot segregation^. An attempt to induce downward flow of liquid during solidification by using a heavy chill plate at the base of the mould was unsuccessful although an increase in the sidewall chill thickness from 30 mm to 50 mm did result in extended columnar crystal growth (Ingot B). This clearly indicated inadequate feeding.

An improved version of Ingot A using heavier chill plates of 80 mm thickness with an insulated feeder head proved considerably more successful as shown in Fig. 5 (Ingot D). Columnar crystal growth further extended towards the core resulting in a relatively narrow central equiaxed segregated zone. The effect of chill plate thickness on columnar crystal growth is shown in Fig. 6, an increase in thickness from 30 mm to 80 mm extending columnar growth by a factor of almost three, indicating the prospect of further potential benefits by an increase in the heat sink. However, weight and handling considerations limited further modification of mould design except for an increase in the volume of the feeder head at the expense of ingot size resulting in mould design E which was used for the casting of all vacuum melted experimental steels. As shown in Fig. 5 the dumbell Ingot C which was placed in chronological casting

' sequence between B and D showed no tendency to segregate indicating lack of directional solidification and inadequate liquid metal feeding.

3.1.3 Macrosegregation in Trial Ingots

Before proceeding with the main programme confirmation of the intensity of segregation was sought by a microanalysis study of Ingots A and D both of which showed the greatest evidence of segregation based on macroetching. Several areas were examined and the analysis was carried out across the longitudinal segregation band using energy dispersive X-ray spectrometry. As shown in Fig. 7 a typical microanalysis scan across the central segregation zone of Ingot A confirmed the existence of semi-continuous spot type segregation over a relatively broad band with a number of manganese peaks reaching values of 1.8-1.9% although the mean value was much lower. At an average bulk manganese content of 1.4% this gave peak and mean segregation ratios of 1.3 and approximately 1.2 respectively. Not surprisingly the segregation band in Ingot D was considerably narrower confirming the expectation based on computed freezing conditions and macroetching results. Segregation was also more intense with higher peak manganese values exceeding 2% and evidence of more continuous band formation.

3.2 Ferri te-Pearli te Steel

As a final check on the suitability of the casting technique a vacuum melt slab ingot was made to a typical X60/X65 low sulphur calcium treated CMnNb linepipe composition, controlled rolled to IS mm plate according to the schedule in Table 2 and subjected to a full microstructural, mechanical properties and HIC assessment. These results were then used as the control to compare the performance of the low carbon PRS and particularly AF/bainite steels which are the main focus of this project..

3.2.1 Microstructure

Representative microstructures together with full plate thickness macrostructure showing the. central segregation band are illustrated in Fig. 8. The low sulphur steel exhibited a typical ferrite-pearlite controlled rolled microstructure. Pronounced pearlite banding was evident in the mid-thickness region of the plate giving way to a more uniform pearlite distribution with discontinuous pearlite colonies and less banding in the surrounding regions. Hard acicular band structures typical of continuously cast material were not pronounced in the sample selected for examination, which was taken from a position close to the top of the ingot, Fig. 3. However, evidence of hard band structures were later found in association with HIC and will be considered in this context.

Although the sulphur content of 0.004% was well in excess of current commercial levels for sour service applications the steel appeared to be relatively clean with uniformly dispersed mainly fully modified globular non-metallic inclusions, having a size distribution less than 2 urn, Fig. 9. No evidence of excessive inclusion segregation could be found in the mid-thickness region although an occasional

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elongated manganese sulphide stringer and massive angular niobium carbonitride particle were evident.

3.2.2 Mechanical Properties

The plate transverse tensile and standard Charpy V notch properties are presented in Table 4 with energy and fracture appearance transition curves in Fig. 10. Strength properties meet the requirements for Grade X65 pipe plate with characteristically high YP/TS ratio of 0.87. Perhaps not surprisingly in view of the relatively low rolling reduction ratio toughness properties were adversely affected to some extent resulting in a 50% FATT (Fracture Appearance Transition Temperature) marginally higher than similar grade commercially produced plate. Nevertheless, the properties represent a close approximation to commercial practice and clearly confirm the establishment of an acceptable laboratory process route.

3.2.3 Standard HIC Assessment

The basic HIC test results are presented in Table 5 together with blister counts and evolved hydrogen contents. Cracking was found in both the BP and NACE solution tests the results reflecting test severity. A visual indication of performance is clearly seen in the C scan traces, Fig. 11, with the NACE solution tests showing the largest projected area of cracking. Considering the BP solution tests the maximum values of CLR per section and per test piece was 14.3% and 8.4% respectively. Of the eighteen faces examined eight contained cracks and the average CLR (including zeros) for all sections was 3.38%. The average number of cracks in the eight sections showing HIC was less than 1.5/section with mean crack length and extent of 1.1 mm and 0.16 mm respectively indicating a moderately BP susceptible steel exhibiting mainly planar type cracking. Relevant crack data for the two test solutions and two plate sampling positions is summarised in Table 6. Comparing the results of the BP solution tests from the two sampling positions it is evident that the .esponse was very similar with no evidence of gross iiihomogeneity between the groups. Four faces from a total of nine examined in each group were cracked giving mean values of crack indices for each group within the expected range of experimental error. Although not necessarily a reliable indicator of HIC performance it is interesting to note that mean evolved hydrogen contents were also similar for each group oftest pieces.

Perhaps surprisingly, considering the greater test severity of NACE solution (pH ~3.5), the proportion of cracked faces was the same as in BP solution. However, the severity of cracking was substantially greater, the total number of cracks increased by a factor of three from 11 to 33 with greater average crack length (1.85 mm) and extent (0.31 mm) values and considerably higher crack indices, the results reflecting a tendency towards longer cracks with a greater stepwise component compared with BP solution tests. As before, the results for the A and B sampling positions were comparable with the mid-length position (Group B) test pieces showing marginally higher crack ratios further confirming the homogeneity of HIC test response and suitability of the experimental techniques.

3.2.4 Relationship Between HIC and Microstructure

The superficial degree of surface blistering as indicated in Table 5 particularly after exposure under NACE solution test conditions in which relatively high hydrogen fugacity is generated at the test piece surface is indicative of intrinsically high HIC resistance.' This was confirmed in an optical metallographic examination of prepared sections from HIC test pieces. Cracking, was restricted to a mid-thickness band of about 1.3 mm maximum width. HIC could not be found in the regions surrounding the centre band which exhibited a relatively uniform ferrite-pearlite microstructure with short discontinuous pearlite bands and absence of gross elongated non-metallic inclusions or inclusion clusters. Centreline HIC tended to be most prevalent along continuous bands of pearlite, pearlite/bainite or bainite/martensite typical of HIC in continuously cast derived flat products, Fig. 12, as reported by other workersM4.i5. Indeed, it has been emphasised that even in super low sulphur calcium treated steel HIC may be found in hard band acicular structures16. From an examination of all the cracked sections HIC rarely developed in the ferrite except during the process of stepwise linkage. On the contrary, arrest of the planar cracks tended to occur at the pearlite/bainite/martensite-ferrite interface. It is also interesting to note that cracks generally occurred at the boundary of the segregation band broadly defining the variable bandwidth as indicated in the photographs of HIC test piece sections, Fig. 13. This behaviour may be

6

explained in terms of the hydrogen concentration profiles during hydrogen charging in relation to the geometry of the HIC sensitive bands35.

3.2.5 Microanalysis of HIC Sensitive Bands

limited microanalysis study of selected areas was carried out to determine segregation profiles in the vicinity of HIC sensitive bands for comparison with typical continuously cast slab products. The area shown in Fig. 12 was chosen for the study which conveniently contained HIC in close proximity within pearlite and bainite/martensite bands. Analyses were made at 5 pm intervals along a number of 10 pm spaced line scans across the segregate bands including cracks as outlined in the electron micrograph included in Fig. 12. The data are available as a series of manganese concentration profiles which may be translated to isometric views of selected area. Although a statistical analysis indicates significant co-segregation of all elements insufficient data are available to include an accurate assessment of. the additional elements.

Typical manganese concentration profiles for the HIC sensitive region ot tne pea r l i t e and bainite/martensite bands are given in Fig. 14 showing peak manganese concentrations of approximately 1.8 and 2.2% giving peak segregation ratios of 1.5 and 1.8 for the pearlite and bainite/martensite ba-Js respectively which is in line with a peak manganese segregation ratio of about 2 in the HIC sensitive acicular bands of continuously cast linepipe steels16,34. These alloy enrichment levels clearly account for the microstructural differences observed confirming that under typical processing conditions in a X60/X65 ferrite-pearlite linepipe steel 1.8% manganese in the segregate band is insufficient to form acicular transformation products but nevertheless may still be susceptible to HIC. An isometric view of manganese concentration in the cracked acicular band is presented in Fig. 15 showing peak concentration at the crack mid-length position decaying together with crack width towards the ends of the band. The results of the microanalysis study further reinforce the suitability of the experimental steelmaking and ingot casting techniques for producing basically high quality HIC resistant material with a central segregation band HIC sensitive microstructure.

3.3 Low Carbon Pearlite Reduced Steel

As discussed previously the low carbon variants of the ferrite-pearlite steels have found favour as sour service linepipe materials with reduced tendency to multiple pearlite banding and HIC sensitive continuous hard band formation. Unfortunately, the strength is limited to Grade X60-X65 and more intense low temperature rolling may be required to compensate the lower carbon content. This has been shown to have an adverse effect on HIC resistance in conventional ferrite-pearlite unmodified and sulphur modified steels35^6 . Since the PRS compositions occupy a natural evolutionary step in the progression towards enhanced HIC resistance their performance will be examined and compared with their predecessors the traditional ferrite-pearlite type materials and their proposed successors the ultra-low carbon AF/bainitic steels. For various reasons it was necessary to make several casts to PRS compositions and the rationale for this will be discussed below.

3.3.1 PRS Composition

The selected composition for the first cast of PRS, ingot/Plate F, was based on commercial practice. Both carbon and manganese contents were reduced from approximately 0.1% and 1.2% respectively in the traditional ferrite-pearlite steel to 0.04% and 1.1%. A vanadium addition was also made to compensate the reduction in carbon and manganese contents. Unfortunately, the cast carbon content was considerably above the aim level giving a composition between ferrite-pearlite, ingot/Plate E, and the intended PRS composition, Table 7. Nevertheless, this was considered to be a useful link between the various steels and the ingot was therefore processed and the properties and HIC performance determined. The intensity of segregation in this plate was subsequently also found to be limited and these two factors necessitated a remake of the PRS composition. This was poured at higher superheat to encourage segregation but while compositional control was good metal spillage unfortunately resulted in a short cast with the possibility of inadequate feeding. A further attempt proved successful in terms of compositional control (ingot/Plate L), and a full properties and HIC assessment was carried out on all plates.

7-

3.3.2 Mechanical Properties

As before transverse tensile properties and Charpy V notch transition curves were determined for each steel. The tensile and basic notch toughness data are summarised in Table 8 with Charpy transition curves in Fig. 16 and typical tensile load-extension curve in Fig. 17. Not surprisingly in view of the relatively high carbon content, Plate F exhibited similar strength and ductility properties to the conventional ferrite-pearlite linepipe steel, Plate E, discussed previously, satisfying Grade X65 strength requirements. However, the Charpy energy and fracture appearance curves showed considerably higher transition temperatures and this again may be attributed in part to the relatively low total rolling reduction ratio and to the additional precipitation strengthening contribution from vanadium.

Considering Plate H from the compositionally acceptable but short cast ingot, tensile properties were very similar to the ferrite-pearlite (Plate E) and high carbon content PRS cast (Plate F) being adequate for Grade X65 pipe plate. A load-extension curve for Plate H is reproduced in Fig. 17 illustrating the characteristic discontinuous yielding behaviour of the PRS and higher carbon family of pearlitic steels. Compared with Plates F and H, Plate L displayed lower flow stress properties by 20-30 N/mm2 reflecting the lower carbon content of 0.03%.

The notch toughness properties of Plates H and L showed remarkable similarity with high ductile shelf energy of ~260 J and low 54 J transition temperature and FATT; a similar order of magnitude to the ferrite-pearlite steel Plate E, and substantially better than the PRS Plate F.

3.3.3 Microstructures

In common with all materials in this vacuum melted series the PRS plates were relatively clean showing a well dispersed distribution of fine fully modified non-metallic inclusions typified in Fig. 18. However, as frequently occurs in HIC studies evidence of gross inclusions was later found in association with HIC.

Photomicrographs illustrating the general microstructural features are shown in Fig. 19. In contrast with the ferrite-pearlite steel, Plate E, the 0.07% carbon PRS Plate F contained relatively well dispersed pearlite colonies with discontinuous pearlite banding towards the plate mid-thickness region. Little evidence of segregate band associated acicular transformation product could be found which is perhaps surprising considering the compositional similarity between this and the ferrite-pearlite Plate E. Hence, the use of higher superheat in an attempt to induce more intense segregation in subsequent melts.

With a reduction in carbon content to the aim level of 0.04%, Plate H, a substantial reduction in the fraction of pearlite was evident accompanied by a relatively coarse deformed ferrite grain structure. A tendency towards centreline segregation was also detected, manifest as discontinuous pearlite banding.

3.3.4 Hydrogen Induced Cracking

Individual HIC test results for each section examined together with blister counts and evolved hydrogen contents are given in Table 9. This includes standard 96 h exposure tests in BP and NACE solution and extended exposure tests in NACE solution for Plate F. The BP solution and NACE extended exposure tests were discontinued in the later casts of pearlite reduced steels. All steels showed limited HIC to some degree as confirmed in the C scan traces for Plates H and L, Fig. 20, and the reported crack ratios, Table 9. However, the results indicate a high level of HIC resistance and minimal surface blistering, Fig. 21. Considering Plate F only one section from eighteen in the 96 h BP solution test and three from eighteen in the NACE solution test were cracked and in each case HIC was associated with surface blisters. In view of the failure to induce centreline HIC in the 96 h tests two further series of HIC test pieces were taken from the sample material remaining from Plate F corresponding to the ingot middle and bottom positions. These were coded C and D for the two positions and HIC tests were carried out in NACE solution for times of 96 h and 192 h. Test piece sectioning for crack counting was. based on C scan assessment.

. Paradoxically, all sections were crack free after the 192 h exposure although the parallel 96 h test on adjacent samples resulted in two cracked sections from eighteen, with a total of three crack systems showing evidence of HIC in the mid-thickness region, Fig. 22;

8

PRS Plates H and L with carbon contents closer to the aim level compared with Plate F showed a greater degree of cracking in the HIC tests. Of the eighteen sections examined for the two steels five were cracked containing a total of eleven crack systems, giving an average CLR value including zeros of 2.73% or 9.81% excluding zeros. The maximum crack length was 4 mm with a maximum CLR per individual section of 22.14%.

3.3.5 Microstructural Aspects

An optical microstructural examination was carried out to determine the important HIC related features. The three crack systems in the mid-thickness region of Plate F after a 96 h repeat test in NACE solution showed a tendency to be associated with isolated areas of carbide containing transformation products which were generally too narrow for reliable microhardness determination. For Plates H and L no evidence of segregation band related transformation products could be found in the crack paths the cracks tending to be random. However, in many cases gross inclusions were found within the cracks. Examples of the various crack related features are shown in Fig. 23.

3.4 Acicular Ferrite/Bainite Steels

A number of AF/bainite steels were made, the aim being to maintain a low carbon content to minimise the potential detrimental effect on centreline segregate band microstructure and hardness and consequently on HIC resistance.

3.4.1 Compositions

Numerous AF/bainite compositional variations are available based on ultra-low carbon and high manganese with various hardenability, dispersion and solid solution strengthening additions. Essentially two compositional types were chosen involving either a molybdenum addition or additions of titanium and boron, the latter combination at two levels of manganese. The aim and melt analyses are given in Table 10. 50 kg vacuum melts were used throughout with low sulphur melting base and calcium treatment. As in the case of the PRS material a remelt of the molybdenum containing AF steel (Plate I), was made at a higher superheat to encourage segregation. Processing to 15 mm plate was nominally identical to all previous material and according to the rolling schedule in Table 2.

3.4.2 Mechanical Properties

Duplicate transverse tensile tests were carried out and standard Charpy V notch energy and fracture appearance transition curves determined for each steel. Basic data are summarised in Table 11 with Charpy transition curves in Fig. 24 and tensile load-extension curves in Fig. 25. The tensile characteristics of the AF/bainite steels tended to be somewhat different from the pearlitic steels and generally higher strength. Considering the molybdenum containing steels, Plates G and I, which are both nominally of identical composition, the yield stress and tensile strength values are 20-30 N/mm2 higher than the PRS plates reaching Grade X75 in Plate I, adequate for X70 pipe plate.. Tensile ductility was also comparable with the low carbon PRS materials. Although not showing full continuous yielding behaviour it is interesting to note the elimination of the yield point drop in Plate G (Table 11), an effect which was less pronounced in the marginally higher yield strength Plate I, Fig. 25(a). Nevertheless, a high YP:TS ratio in the region of 0.8-0.9 was comparable with controlled rolled pearlitic steels.

Both the low manganese (Plate K) and high manganese (Plate I) variants of the CMnTiB steel displayed continuous yielding behaviour giving a substantial increase in tensile strength of over 100 N/mm2 and a decrease in YP:TS ratio compared with the pearlitic steels. However, the most useful flow characteristics were only developed in the C 1.9% MnTiB steel, Plate K, which displayed a high level of strain hardening with YP:TS ratio of 0.70 and a 0.2% PS comparable with the pearlitic steels. Based on the stress at 2% elongation criterion the strength properties of this steel approach Grade X80.

All the AF/bainite steels showed a high level of notch toughness with ductile shelf energies ranging from 175 to 270 J and 54 J transition temperatures of -64 to -97G. Except for the CMnMoNb Steel G, which had a relatively high FATT, the 50% FATT of the remaining steels was low comparable with the pearlitic

- 9

steel, Plate E, and substantially better than the PRS plates. However, in terms of overall mechanical property considerations the high manganese CMnTiB steel, Plate J, exhibited potentially the most promising balance of properties.

3.4.3 Microstructures

Representative microstructures of the three steel types are illustrated in Fig. 26 and not surprisingly considering the low carbon content are characterised by a low fraction of free carbide which tends to be highest in Plates I and K the less hardenable of the AF/bainite compositions. The CMnMoNb steel most closely resembles the microstructure of the PRS materials in having a mainly polygonal ferrite microstructure with a small fraction of elongated carbide phase, in this case comprising pearlite and bainite. Both of the CMnTiB steels exhibit highly dislocated fine acicular microstructures with, in the case of Plate J, a high degree of homogeneity and little evidence of second phase. The relatively uniformly dispersed carbide phase identified as upper bainite in Plate K contrasts with the discreet islands of second phase in the CMnMoNb steels.

Little evidence of major centreline segregation could be found by optical microscopy, partly as a consequence of the low carbon content and the effect on etching characteristics and also perhaps attributable to a reduced level of segregation in the low carbon compositions. However, minor centreline banding and microstructural differences between mid-plate thickness and surrounding regions may be observed in the accompanying micrographs.

3*4.4 Hydrogen Induced Cracking

The crack indices CLR, CTR and CSR for standard and extended exposure HIC tests with blister assessment and evolved hydrogen contents are presented in Table 12. Because Plate G was the first of the AF/bainite steels to be produced tests were carried out in both BP and NACE solution to assess the general response of this category of materials. No cracking could be found in the standard 96 h tests and consequently a further series of 96 h and 192 h tests were conducted in NACE solution on test pieces from the equivalent ingot mid-length and bottom positions of the plate and coded C and D respectively. Test piece sectioning was based on C scan assessment. While marginally more susceptible to HIC the severity of cracking was relatively low with only one section cracked in the 96 h test and two sections in the 192 h test giving a total of three cracked sections from 36 sections examined. From the photograph of HIC test piece sections in Fig. 27 two of the cracks in Plate G may be identified as centreline HIC. As previously discussed in relation to the AF/bainite steels all subsequent melts of AF steels following Steel G were cast at higher superheat to help induce segregation.

NACE solution HIC tests on the remaining AF/bainite steels I, J and K which were subjected to though not sectioned on the basis of C scan, Fig. 28, showed a greater sensitivity than both the PRS and the initial CMnMoNb AF/bainite Steel G. A similar analysis of the data revealed fewer crack systems from a greater proportion of cracked faces and longer crack lengths resulting in substantially higher values of CLR. The maximum CLR value per section was also higher at 46% as was the CTR arising from a greater stepwise component of HIC compared with the low carbon pearlitic steels. This suggests a greater hydrogen sensitivity of the AF/bainite steels which tended to show greater maximum crack length and higher mean CLP excluding zeros, the higher the tensile strength.

In accordance with standard practice crack systems lying entirely within 1 mm of the test piece upper and lower surfaces were not included in the ratio calculations. A large proportion of these subsurface cracks are manifest as surface blisters and as seen in Fig. 29 the tendency was towards a greater severity of surface blistering in the AF/bainite steels compared with the ferrite-pearlite and PRS compositions. Evolved hydrogen content was also sensitive to steel type, increasing from approximately 1 ml/100 g in the pearlite reduced steels to over 3 ml/100 g in the highest tensile strength AF/bainite steel, Plate J.

3.4.5 Microstructural Aspects

Unlike the well defined central segregation band cracking in the 0.1% C ferrite-pearlite steel, Plate E, the AF/bainite steels behaved similarly to the PRS plates in displaying essentially random HIC behaviour as

- 1 0 -

illustrated in Fig. 27.- Sections were etched in nital and subjected to an optical microstructural examination to determine the presence of crack related features and particularly macrosegregation. As shown in Fig. 30 segregation band transformation products were much less obvious in these low carbon steels compared with the 0.1% C ferrite-pearlite steel. Gross inclusions were, however, found in some cracks. It must therefore be concluded that on the basis of microstructural examination little evidence could be found for the existence of gross macrosegregation in the AF/bainite steels and consequently is unlikely to be implicated in the apparent hydrogen sensitivity.

3.5 Effect of Phosphorus

Together with manganese, phosphorus has been indicted as an undesirable element in the segregation band giving segregation ratios of about 10, and resulting in substantial HIC19. Although not entirely proven it is claimed that the potent hardenability effect of phosphorus3^ leading to hard band acicular transformation structures is the principal cause of cracking and very low phosphorus content is now being advocated19. In view of the commercial implications a limited study of the effect of phosphorus has been included in the test programme.

3.5.1 Compositions and Mechanical Properties .

The steel types selected were those showing the greatest HIC sensitivity in the current test programme, i.e. the 0.1% C ferrite-pearlite and 1.9% Mn CMnTiB AF/bainite steels. The levels of phosphorus chosen were 0.005% for the ferrite-pearlite steel with 0.003% and 0.020% for the two CMnTiB steels compared with the 0.011-0.015% phosphorus content of all previous materials, Table 13. Unfortunately, the carbon and titanium contents of the low phosphoi is AF/bainite steel, Ingot N, were substantially below the aim levels making a comparison of the effect of phosphorus difficult. Nevertheless, general trends will be apparent.

Tensile and notch toughness properties for the three plates are given in Table 14 with Charpy transition curves and tensile load-extension curves in Figs. 31 and 32 respectively. Although showing marginally lower strength the tensile properties of the low phosphorus ferrite-pearlite steel, Plate M compare favourably with the higher phosphorus variant of this material, Plate E, reported previously, satisfying Grade X65 requirements. The differences in strength properties between the two casts amounting to 18 N/mm2 and 25 N/mm2 for yield stress and tensile strength respectively may largely be accounted for by the solid solution strengthening effect of phosphorus and minor deviations in rolling temperature.

A comparison of tensile properties for the AF steels, Plates N and O, which were intended to be nominally identical, reveal major differences which appear to be directly attributable to compositional effects. As explained earlier the carbon and titanium contents of Plate N were substantially below the aim levels. This factor coupled with the intended lower phosphorus content has presumably combined to reduce hardenability, resulting in discontinuous'type yielding behaviour with X65 strength properties almost identical to the pearlitic steel though with markedly higher ductility properties. As expected, the high phosphorus AF steel, Plate O, displayed continuous yielding characteristics with low flow stress and high strain hardening rate. Based on the criterion of stress at 2% extension the yield stress was marginally above X80. The tensile ductility, however, appeared to be inferior compared with Plates M and N and indeed with earlier casts of AF steels, although the abnormal tensile fractures (see Table 14) may at least partially account for the difference.

A high level of notch toughness was displayed by all steels, comparable to earlier casts of similar compositions with ductile shelf energies well in excess of 200 J and low transition temperatures. Notwithstanding the compositional differences between the two AF steels the toughness properties were remarkably similar with marginally lower ductile shelf energy and transition temperature in the high phosphorus alloy.

3.5.2 Hydrogen Induced Cracking

In order to ensure a statistically significant result from the HIC tests and to maximise the chances of including segregated material sampling was carried out at three adjacent positions along the plate length

11

as shown in Fig. 33. Standard HIC test data for the three steels including crack indices, blister counts and evolved hydrogen contents are presented in Tables 15-17. Except for major cracks in one section from each of test pieces B3 and B5 giving individual CLR values of 9% and 44%, hydrogen damage was superficial in the low phosphorus ferrite-pearlite steel, Plate M. This was substantiated in the C scan traces, a selection of which are presented in Fig. 34 for illustrative purposes. From a total of 27 sections examined 18 were clear and 7 contained minor subsurface cracks of lengths 0.1-0.2 mm. Only in the two severely cracked sections, B3C and B5A, could central segregation band HIC be identified with crack lengths of 1.8 and 7.6 mm. Surface blistering was also superficial as seen in Fig. 35.

An analysis of crack data for all steels including the phosphorus variants is summarised in Table 18 and Fig. 36. From this the substantial effect of phosphorus on HIC susceptibility of Grade X65 controlled rolled ferrite-pearlite steel will be evident. Although a reduction in phosphorus content from 0.015% (Plate E) to 0.005% (Plate M) had a relatively minor effect in proportionately reducing the number of crack systems the severity of cracking was noticeably decreased. Mean crack length was reduced from 1.85 to 0.85 mm with a similar proportionate reduction in crack extent. The net result was a decrease in CLR from 17% to 2.2% and a decrease in CTR from 4% to 0.5%, almost entirely eliminating the stepwise

. component of HIC.

In view of the unintended differences in carbon content and tensile properties between Plates N and O a strict comparison of the effect of phosphorus is precluded. Nevertheless, it is interesting to observe a pronounced difference in HIC response between the two casts the low phosphorus low carbon steel, Plate N, displaying substantially the higher degree of HIC resistance, Tables 16 and 17. From a total of 27 sections-examined 22 were clear and five contained cracks with a total of seven crack systems and a maximum CLR per test piece of 16.5%. Only one section C3C contained plate mid-thickness cracks.

In contrast the high phosphorus steel, Plate O, had a total of 11 clear sections from 27 with 37 crack systems and a maximum CLR per test piece of 38%. However, cracking was almost entirely restricted to the surface regions of the test pieces with little identifiable mid-thickness cracking as was the case with the low phosphorus steel. This is reflected in the severity of surface blistering with many major blisters greater than 5 mm diameter as seen in Fig. 35. From an analysis of the HIC data in Table 18 the beneficial effect of reduced HIC sensitivity in the low phosphorus steel is evident, mean CLR and CTR values being reduced from 12% and 3% in the high phosphorus steel to 3% and 1% respectively. Surprisingly the mean crack length and extent values were lower in the high phosphorus steel and contrary to expectations considering the high crack ratios.- This, however, is explainable by the large number of minor cracks effectively reducing the mean values. It is also interesting to note the consistently higher evolved hydrogen values and corrosion weight loss in the medium and high phosphorus 1.9% manganese CMnTiB AF/bainite steels, Plates J and O.

Comparing the overall performance of the,.various microstructure/compositional types with the relatively susceptible ferrite-pearlite control steel, Plate E, it will be observed that the three PRS plates, F, H and L, showed a substantially lower level of susceptibility in terms of mean crack ratios, crack dimensions and number of crack systems. Their performance was only equalled in the ferrite-pearlite steel by a reduction in phosphorus content to 0.005%, Plate M, which is attributable to reduced sensitivity in the segregation band. The AF/bainite steels, however, behaved rather differently from the PRS plates displaying an initially higher level of susceptibility and progressively increasing sensitivity through the compositional series CMnMoNb, Plates G and 1,1.6% Mn CMnTiB, Plate K and finally to the 1.9% Mn CMnTiB Plates O and J. It would seem therefore that while relatively good HIC resistance may be achieved in the low carbon pearlitic steels with a phosphorus content typically of about 0.011% the AF/bainite type materials require lower phosphorus content to achieve a similar level of HIC resistance.

3.6 Single Sided HIC Tests

In order to fully assess the HIC response of the various microstructural and compositional variables proposed it was considered necessary to carry out long term single sided exposure tests and to monitor crack initiation and growth together with hydrogen permeation simultaneously during the exposure period. A multi test piece cell was therefore constructed capable of testing plate or unfiattened pipe samples having dimensions approximately 150 mm x 150 mm with an exposed area per test piece of

- 1 2

100 cm2, Fig. 37. Test solution was contained in a temperature controlled reservoir, saturated with H2S and circulated within the test loop, Fig. 38, at a flow rate of approximately 1 litre/min.

3.6.1 Commissioning Trial

Prior to commencement of the single sided testing programme a commissioning trial was carried out to highlight any problems and to establish a working procedure. The trial was conducted on a HIC susceptible linepipe steel using NACE solution formulated according to TM-01-77, since this solution was expected to show greater deterioration with respect to acidity compared with BP solution. A solution volume to test piece surface area ratio of 30 ml/cm? used and regular monitoring of pH and H2S concentration was carried out throughout the 1000 h test. The test solution temperature was measured in the reservoir vessel and also in the return leg of the test cell.

Crack development was monitored using an ultrasonic C scan technique with the test pieces in situ on the test cell. A jig for this operation is shown in Fig. 39, in which the test cell remains stationary for the scanning of plane surfaces or rotates at a predetermined radius of curvature for the scanning of unflattened pipe samples. Although not carried out in the initial trial hydrogen permeation was determined in a subsequent test using a Hellesens mini cell. After removal from the chemical cell the test pieces were examined for general surface condition of the exposed faces and through thickness sections across the rolling plane were metallographically prepared for examination of cracks. A limited manganese analysis was also carried out by energy dispersive X-ray spectrometry in suspected segregate regions of cracks. To avoid unnecessary duplication these aspects oftest procedure will not be discussed with respect to the commissioning trial but will be considered within the context of the main test programme.

3.6.2 Operational Performance

As shown in Fig. 40 the pH of the test solution rapidly increased during the first 48 h of exposure from 2.7 to about 3.7 and thereafter remained virtually unchanged throughout the duration of the test. Thus the pH remains below the frequently specified maximum of 4.0 but more important the test is conducted at essentially constant pH, unlike the 96 h total immersion HIC test in which pH changes over a major period of the test. Similarly the H2S concentration rapidly stabilised at approximately 2700 ppm which is within the 2300-3500 ppm concentration range for saturation. A test solution volume ratio of 30 ml/cm2 therefore appears to be adequate to avoid deterioration over a 1000 h test duration.

Temperature control of solution in the reservoir was very stable and better than 25 0.25C although temperature on exit from the test cell was consistently lower than the reservoir temperature by 1-1.5C which is within the specified range of 25 3C for HIC testing. Nevertheless if desired this temperature drop can be compensated by an increase in the reservoir temperature.

Frequent ultrasonic checks were carried out to determine the time for crack initiation and the rate of crack growth. A typical series of C scan traces showing crack development and the resultant crack area plot are shown in Figs. 41 and 42 respectively. Once initiated in the early stages of testing cracks frequently grow at a higher rate along the rolling direction and at the expense of new crack systems. Also in plate of continuously cast origin major cracking occurs in the centreline segregation region.

3.6.3 Single Sided HIC Test Response of Experimental Steels

In order to realistically assess the performance of the experimental materials single sided tests were carried out on a representative sample of the ferrite-pearlite, PRS and AF/bainite steel type, Plates E, I, J, K and L. In all cases test pieces were linished to a 320 grit finish at a thickness of 13.3 mm. The HIC susceptible ferrite-pearlite steel, Plate E, was tested in both BP and NACE solution while the remaining steels were tested in NACE solution only.

Considering the response of Plate E, cracking did not occur in BP solution and only a relatively small total area of HIC from three crack systems could be detected when exposed in NACE solution, Fig. 43. These cracks representing approximately 0.9% of the exposed area were detected within the first 70 h of

13

test and remained essentially stable thereafter. Hydrogen permeation studies indicated hydrogen breakthrough times of about 5 h and 9 h for NACE and BP solutions respectively. As would be expected the rise time was faster and the equilibrium hydrogen concentration higher under NACE solution test conditions compared with BP solution, Fig. 44. Using this and additional data from hydrogen charging experiments and standard HIC tests, preliminary estimates of hydrogen concentration at the centre of Plate E have been made under single sided exposure conditions and compared with values for commercial HIC susceptible linepipe steels under similar test conditions, Fig. 45. Computed hydrogen diffusion coefficients and critical hydrogen concentrations are given in Table 19. The method of calculation is given in the Appendix.

Comparison of the diffusion coefficients for the commercial linepipe steels and the vacuum melted calcium treated experimental steel, Plate E, shows that the latter has a greater coefficient by one order of magnitude which is in line with the view that the presence of traps greatly reduces transport of hydrogen through the lattice. Hence a steel with a high diffusion coefficient will reach the critical hydrogen concentration in the centreline segregation region more quickly than one with a low coefficient. Since the early stages of the single sided test on Plate E were not ultrasonically monitored, coinciding with a weekend, it is difficult to precisely define the critical hydrogen concentration for HIC initiation from Fig. 45 unlike the data for the commercial steels, 1 and 2, which indicate values in the region of 1 ml/100 g.

The cracked single sided HIC test piece tested in NACE solution was sectioned across the crack areas as defined by the ultrasonic C scan assessment, Fig. 43, and examined for microstructural related evidence of crack nucleation and growth. Clearly the cracks appeared to be nucleated at inclusion cluster sites followed by linkage. It is interesting to note that the crack geometry did not coincide with segregate zone martensite/bainite or pearlite bands but propagated principally within the ferrite as seen in Fig. 46 suggesting inclusion dominated growth.

Unlike the ferrite-pearlite steel HIC was not detected in the PRS and AF/bainite steels during the 1000 h exposure despite peak hydrogen permeation concentrations, Fig. 47, which were approximately equivalent to and in two cases (Plates I and K) marginally exceeded the level achieved in the superficially hydrogen damaged Plate E. The permeation curves for all the low carbon materials, unlike the ferrite-pearlite steel also showed pseudoclassical permeation behaviour with rapid current rise followed by a decay which has been attributed to film formation^. The reason for these differences remain obscure although it would appear not to be a contributory factor in the cracking of Plate E since damage occurred in the early stages oftest and did not develop despite prolonged exposure under peak hydrogen conditions. A further interesting feature is the substantially lower permeation current for the PRS.Plate L compared with the ferrite-pearlite and AF/bainite steels. Significantly PRS are the only materials with a vanadium addition suggesting carbonitride precipitates as high energy hydrogen trap sites, a claim which has been made with regard to titanium4".

3.7 Weldability Studies

Laboratory simulated linepipe submerged arc seam welds were made on five steels covering the full range of pearlitic and AF steels investigated. Chemical compositions of the various steel types are given in Tables 1, 7 and 10. Since the quantity of plate was limited only one weld was attempted for each steel using the best choice of weld consumables compatible with plate chemistry. Oerlikon basic agglomerated OP122 flux was employed in all cases together with 4 mm diameter Tibor 33 wires for the ferrite-pearlite and PRS Plates E and L and Tibor 22 for the various AF steels, Plates I, J and K. Plates of nominally 15 mm thickness and 200 mm wide were slit along the rolling direction at the mid-width position. Double V preparation with 5.5 mm nose was machined in the 600 mm long plates using 50 and 40 outside and inside angles respectively. The plates were batched in multiples of three and two as determined by weld consumables and arranged as illustrated in Fig. 48 with scrap plate heat sinks. A continuous tacking weld was made in the root of the outside preparation using the CO2 GSM A welding process with 1.2 mm diameter Murex Bostrand LW1 wire. Triple wire submerged arc single pass" welds were made with a heat input of 3.22 kJ/mm for the inside and 4.58 kJ/mm for the outside. Welding parameters were monitored at 1 s intervals using a calibrated weld monitor with digital printout. Details of the welding procedure are given in Table 20. Notch toughness of the weld metal was determined with the V notch corresponding

14

to the weld metal centreline. HIC tests were also carried out using standard sampling procedure with the test piece axis across the weld line and transverse to the plate rolling direction.

3.7.1 Notch Toughness of Submerged Arc Welds

Notch toughness data for the submerged arc simulated seam welds is given in Table 21 and Fig. 49. In all steels the notch toughness appeared to be compatible with plate properties having high shelf energies and low transition temperatures well within current recognised limits for linepipe seam welds. Taking a recent stringent toughness specification of 50 J at -30C it will be evident that a large toughness margin exists both for the traditional ferrite-pearlite steels and the more highly alloyed AF compositions.

3.7.2 HIC Tests on Welds

The results of standard HIC tests on simulated seam weld test pieces of pearlitic and F steels are summarised in Tables 22 and 23. Although substantial mid-thickness segregate band cracking occurred in the parent plate sections of HIC test pieces (Sections A and C) from the ferrite-pearlite steel, Plate E, thereby corroborating earlier results, the weld metal containing sections (Section B) generally had fewer cracks. From an examination of the weld metal containing sections (B sections) only one crack was found in association with the weld, apparently nucleating in the HAZ and arresting in the weld metal. In the case of the pearlite reduced steel, Plate L, the general level of cracking was considerably reduced compared with Plate E in agreement with previous finding. Only one weld metal containing section showed HIC and this was well clear of the weld region.

Considering the AF steels, Plates I, J and K, no evidence of weld metal HIC could be found, indeed only one B section in the 1.9% Mn CMnTiB AF steel, Plate J, contained a single crack in the parent microstructure/HAZ region of the test piece. As in the case of the ferrite-pearlite steels the level of cracking in the parent plate portion of the test pieces was generally in agreement with earlier reported findings. Typical surface blister appearance of all steels is illustrated in Fig. 50 and again corroborates previous results. The results of the welding study therefore demonstrate that the low carbon enhanced alloy AF/bainite steels may be laboratory submerged arc welded to simulate SAW pipe practice with no apparent detrimental effect on HIC performance.

3.8 Microanalysis and Microhardness

To complete the investigation a microanalysis and microhardness study was conducted on selected HIC test pieces from the PRS, AF/bainite and low phosphorus ferrite-pearlite steels and also on the cracked single sided ferrite-pearlite steel testpiece, Plate E, to confirm the relative importance of segregation band composition and hardness in relation to HIC. The microanalysis study was limited principally to manganese determination.

Considering compositional aspects of HIC the crack system on the single sided test piece from Plate E shown in Fig. 46 which showed evidence of inclusion dominated HIC was chosen for study. Two areas were selected, in the immediate vicinity of the branched crack system, Fig, 46(b), and in an uncracked multisegregate band region. Typical manganese concentration profiles for various positions along the crack system are reproduced in Fig. 51 showing peak values of less than 1.4% for a base manganese content of 1.23%, confirming the view gained from the metallographic study that the cracks propagated in basically unsegregated regions of the matrix and were essentially inclusion activated. Further confirmation of this view was provided by the high aluminium content detected within the crack, suggesting alumina rich non-metallic inclusions. In contrast, the uncracked region of a segregate band contained relatively high manganese concentration peaks occasionally exceeding 1.9% in small areas but usually within the range 1.8-1.9% as seen in the contour map, Fig. 52. Evidently this concentration is too low for the nucleation and growth of HIC in the absence of gross inclusions under single sided exposure conditions, and tends to support the findings from standard HIC tests reported previously that in the relatively clean 0.1% CMnNb steel a segregation band manganese concentration greater than about 2% is required for crack propagation.

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A standard HIC testpiece from the low phosphorus variant of the ferrite-pearlite steel, Plate M, was subjected to a microanalysis study across one of the two mid-plate thickness cracks found in this material. The results of an EDAX analysis across this crack, Fig. 53, revealed peak manganese concentrations of approximately 1.4% to 1.6%, depending on the interpretation of the concentration profile, corresponding to a manganese segregation ratio range of 1.13-1.29. Clearly this is well below the peak ratios of 1.5 and 1.8 for the HIC sensitive pearlite and bainite/martensite regions in the higher phosphorus ferrite-pearlite steel, Plate E, as indicated previously. A further surprising factor in the work on low phosphorus ferrite-pearlite steel is the apparent lack of co-segregation of other elements.

As demonstrated earlier segregation was difficult to detect in the low carbon steels, particularly in the AF/bainite types, using standard optical metallography. This necessitated a limited microanalysis study to determine the degree of substitutional element segregation. With regard to the PRS compositions, Plates F and L were chosen representing high and marginally low carbon variants of this material. The region of Plate F chosen for analysis was in the vicinity of a crack which had propagated along a segregate band and become stepwise leaving a clearly visible uncracked extension of the hard band microstructure. A similar feature was also selected for Plate L. These areas are illustrated in the photomicrographs, Fig. 23. Microanalysis data are presented in Fig. 54. The line scans across the crack system and segregate band in Plate F show a manganese concentration of about 1.8% in the crack region decreasing to less than 1.6% in the uncracked area giving manganese segregation ratios of 1.7 and 1.5 respectively. This further confirms the powerful effect of segregation band hardenabil i ty element concentration on HIC propagation. It is also interesting to note the narrowing of the segregation band in the uncracked extension of the HIC which may have been a further factor in restricting crack propagation.

Considering the low carbon PRS, Plate L, the maximum manganese concentration found in the vicinity of the centreline HIC was 1.2-1.3% giving a low concentration ratio less than 1.2. The major part of the crack apparently propagated in the unsegregated 1.1% Mn matrix tending to confirm the view that HIC was principally inclusion dominated. Nevertheless, it is interesting to note that a small amount of hydrogen damage appears to have nucleated wholly within a 1.2-1.3% Mn band as seen in Fig. 54(b).

A similar analysis to that described above was also conducted on the AF/bainite steels, Plates G, J and K, and subsequently on Plates N and O the low and high phosphorus variants of Plate J. A contour map of manganese concentration was constructed for the branched crack region of the CMnMoNb steel, Plate G, illustrated in photomicrograph Fig. 30(b). This shows crack propagation in segregation bands containing 2% manganese or greater and apparent arrest in lower manganese areas of the microstructure, Fig. 55. Surprisingly the peak concentration of manganese detected was of the order of 2.2% giving a relatively low segregation ratio of 1.2. These findings were also confirmed in a study of the stepped crack system illustrated in Fig. 30(a) showing the main and intermediate crack system confined principally to the high manganese regions of the microstructure, Fig. 56.

The CMnTiB AF/bainite steels, Plates J and K, showed relatively low levels of manganese segregation and little apparent preference of the crack path with respect to compositional variables, Fig. 57. In both cases the areas of maximum manganese concentration ratio through which cracks propagated was less than 1.2 giving maximum concentrations of 2.1% and 1.8% for Plates J and K respectively.

Final confirmation of the generally low segregation ratio profiles in the low carbon steels was obtained by microanalysis across a mid-thickness band covering a distance of about 4000 pm. The results for Plates J and K, Fig. 58, clearly show relatively low peak manganese concentrations of about the same order as those obtained in Fig. 57, indicating a strong dependency of solute segregation on carbon content in agreement with published data41. Lack of co-segregation of the solute elements both at the plate centreline and in the vicinity of HIC was also confirmed in the low and high phosphorus AF/bainite steels, Plates N and O, Fig. 59.

An examination of microhardhess in the segregate band region of HIC test pieces from a representative cross section of the various microstructural types revealed a pattern of behaviour similar to that shown by phosphorus. The severity of cracking as determined by CLR was dependent on steel type, Fig. 60. Within the range of data points for hardnesses of about 210-250 HV the most HIC sensitive steels appear to be the AF/bainite types with mean CLR values of approximately 10%. Significantly the pearlite reduced steels

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and the low phosphorus ferrite-pearlite steel with equivalent segregation/HIC band hardnesses showed low values of CLR as did the low phosphorus, low carbon AF/bainite steel which had a mean segregation band hardness of 187 HV. Not surprisingly the medium phosphorus ferrite-pearlite steel showed the highest hardness and CLR values. While individual values of microhardness above 300 HV were occasionally found it is interesting to note that this was rarely equalled even in cracked pearlite/bainite bands contrary to the general belief that 300 HV is the critical hardness1^ above which significant cracking occurs.

4. CONCLUSIONS

A laboratory slab casting technique has been developed capable of simulating centreline segregation typical of continuously cast products.

Using this technique a range of calcium treated low sulphur vacuum melted steels covering the microstructural types ferrite-pearlite, pearlite reduced and low carbon acicular ferrite/bainite have been produced and controlled rolled to 15 mm thick plate. HIC tests on these materials in H2S saturated NACE solution have shown the 0.1% CMnNb ferrite-pearlite Grade X65 steel to be the most sensitive with a relatively high value of CLR and pronounced cracking in the martensite/bainite and pearlite central segregation band. Manganese segregation peaks of approximately 1.8% and 2.2% were found giving segregation ratios of 1.5 and 1.8 for pearlite and martensite/bainite bands respectively. The low carbon X60-X65 CMnNbV pearlite reduced steels were the least susceptible to HIC and gave mean CLR values less than 3% with limited central segregation band cracking. The maximum manganese segregation ratios found in the vicinity of HIC were 1.7 and 1.2 for steels containing 0.07% and 0.03% carbon content respectively.

HIC susceptibility intermediate between that for the ferrite-pearlite and PRS steels was found for the X70-X80 CMnMoNb and CMnTiB AF/bainite steels with cracking tending to be random as in the PRS plate. The maximum manganese segregation ratio was also about 1.2 and comparable with low carbon PRS.

Surface blistering was more pronounced in the AF/bainite steels and particularly the C1.9MnTiB steel compared with the PRS and ferrite-pearlite steels.

All steels were highly resistant to HIC in the single sided HIC test, only the ferrite-pearlite steel showing minimal crack development. A.reduction in phosphorus content from about 0.011 to 0.003-0.005 in the ferrite-pearlite and C1.9MnTiB AF/bainite steel resulted in a pronounced beneficial effect on HIC resistance reducing mean CLR values from approximately 15% to 3%, comparable with the highly resistant PRS.

Mean segregation band microhardness varied from 187 HV25 g in the low phosphorus C1.9MnTiB AF/bainite HIC resistant steel to 281 HV25 g in the CMnNb ferrite-pearlite susceptible steel. The low susceptibility PRS and low phosphorus CMnNb ferrite-pearlite steels had mean segregation band hardnesses of 227 to 240 HV25 g comparable with the susceptible CMnTiB AF/bainite steels. All steels showed good weldability with a high level of weld metal toughness and HIC resistance.

5. REFERENCES

1. Miyoshi, E., Tanaka, T., Terasaki, F. and Ikeda, A., ASME Publication, Paper No. 75-Pet-2 (1975).

2. Coldren, A.P. and Tither, G., Journal of Metals, 5th May, 1976.

3. Wilde, B.E., Kim, CD. and Phelps, E.H., Int. Conf. Corrosion 80, Paper No. 7.

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4. Taira, T., Tsukada, K., Kobayashi, U., Inagaki, H. a