The Influence of Heat Input on the Fracture and …cradpdf.drdc-rddc.gc.ca/PDFS/unc82/p531314.pdfDefence R&D Canada – Atlantic DEFENCE DÉFENSE & The Influence of Heat Input on the

Defence R&D Canada – Atlantic

DEFENCE DÉFENSE&

The Influence of Heat Input on the Fracture

and Metallurgical Properties of HSLA-65

Steel Welds

Fabrication conditions, Tensile, Impact andMicrostructures

Christopher BayleyDRDC AtlanticDockyard Laboratory Pacific

Adam ManteiUniversity of VictoriaDept. of Mechanical Engineering

Technical Memorandum

DRDC Atlantic TM 2008-130

October 2008

Copy No. _____

Defence Research andDevelopment Canada

Recherche et développementpour la défense Canada

This page intentionally left blank.

The Influence of Heat Input on the Fracture and Metallurgical Properties of HSLA-65 Steel Welds Fabrication conditions, Tensile, Impact and Microstructures

Christopher Bayley DRDC Atlantic Dockyard Laboratory Pacific Adam Mantei University of Victoria Dept. of Mechanical Engineering

Defence R&D Canada – Atlantic Technical Memorandum DRDC Atlantic TM 2008-130 October 2008

Principal Author

Original signed by Christopher Bayley

Christopher Bayley

Defence Scientist

Approved by

Original signed by Terry Foster

Terry Foster

H/DLP

Approved for release by

Original signed by Ron Kuwahara for

James L. Kennedy

Chair DRP

© Her Majesty the Queen in Right of Canada, as represented by the Minister of National Defence, 2008

© Sa Majesté la Reine (en droit du Canada), telle que représentée par le ministre de la Défense nationale, 2008

Abstract ……..

HSLA-65 steel (ASTM A 945-05) is regarded as an excellent naval ship steel. Its low carbon equivalent and strength-to-weight ratio make it a candidate for future Canadian naval platforms. Unlike previous naval platforms, future ships might be expected to be deployed to Arctic waters. This increased operational requirement imposes a strict fracture toughness criterion for both the base metal and welds. In particular, the heat affected zone near the weld is of concern, especially for high heat input processes. In this study, welds were fabricated using the flux core arc welding and Gas Metal Arc processes with variations in the welding heat input and welding consumable. While all of the welded panels exhibited satisfactory tensile and hardness properties, a dramatic shift in the ductile-to-brittle transition temperature to higher temperatures was observed for high heat input conditions when the notch was positioned within the heat affected zone. The interpretation of this shift in transition behaviour was confounded by the relative width of the coarse grained heat affected zone, the microstructure sampled by the advancing crack and the position of the blunt Charpy notch.

Résumé

L’acier du type HSLA-65 (norme ASTM A 945-05) est considéré comme un excellent acier pour les navires militaires. Sa faible teneur en équivalent carbone et le bas rapport résistance-poids qui le caractérisent en font un matériau intéressant pour la construction de futures plates-formes navales de la flotte canadienne. Contrairement aux anciennes plates-formes navales, les navires du futur pourraient être utilisés dans les eaux arctiques. Ce besoin opérationnel supplémentaire impose un critère très rigoureux en matière de résistance à la rupture, pour le métal de base comme pour les soudures. La zone thermiquement affectée située près d’une soudure constitue un sujet de préoccupation particulier, notamment dans le cas de procédés à apport de chaleur élevé. Dans le cadre de l’étude dont les résultats sont fournis dans le présent rapport, on a fabriqué des joints soudés en utilisant deux procédés de soudage distincts, soit le soudage à l’arc avec fil fourré et le soudage à l’arc sous protection gazeuse avec fil plein, en faisant varier l’apport de chaleur lors du soudage et la nature de l’électrode consommable. Bien que tous les panneaux soudés mis à l’essai présentent des propriétés satisfaisantes en matière de résistance à la traction et de dureté, on a observé un déplacement remarquable de la température de transition ductile-fragile vers des valeurs plus élevées, dans des conditions d’apport de chaleur élevé et lorsque l’entaille est effectuée dans la zone thermiquement affectée. La largeur relative de la zone thermiquement affectée à grain grossier, la microstructure de l’échantillon prélevé dans la région de progression de la fissure et la position de l’entaille en V peu aiguë, lors de l’essai de résilience Charpy, ont rendu plus complexe l’interprétation de cette modification du comportement de transition.

DRDC Atlantic TM 2008-130 i


ii DRDC Atlantic TM 2008-130

Executive summary

The Influence of Heat Input on the Fracture and Metallurgical Properties of HSLA-65 Steel Welds: Fabrication conditions, Tensile, Impact and Microstructures

Christopher Bayley; Adam Mantei; DRDC Atlantic TM 2008-130; Defence R&D Canada – Atlantic; October 2008.

Introduction or background: HSLA-65 (ASTM A945 Grade 65) is a highly weldable steel that is characterized by its low carbon and alloy content, and hence carbon equivalent. It was designed to fill the niche between higher strength HSLA-80 pipeline steels and conventional shipyard high-strength steels (i.e. American Bureau of Shipping DH/EH-36). HSLA65 was developed to have improved strength, toughness and weldability for naval ship structures. These steels depend on low carbon content and microalloying additions, in conjunction with thermo-mechanical processing, to obtain their properties, rather than relatively expensive alloying and/or quench and tempering. Design studies of destroyer-type hulls showed that the strength benefits of 80 ksi steels were limited by plate buckling and that a 65 ksi yield strength steel could achieve similar weight. Furthermore, fabrication and material costs of HSLA-65 are similar to shipyard high-strength steels.

Results: The report summarizes the welding conditions, tensile strength, hardness and impact energy for five different welded panels constructed using either the gas metal arc (GMAW) or flux core arc (FCAW) welding processes. For the FCAW panels, the influence of heat input and weld consumable are investigated and compared with the GMAW welds fabricated using an intermediate strength consumable and heat input level.

While the tensile strength and hardness of all the welding panels are satisfactory there was a significant shift in the ductile-to-brittle transition behaviour in the heat affected zone (HAZ) of the high heat input welds. This drastic shift in the transition behaviour was found to a function of the crack propagation within the microstructural gradient of the HAZ. The propagation of the crack through different microstructural regions also accounts for the increased scatter in the ductile-to-brittle transition behaviour of the high heat input welds.

Significance: These results highlight the need for stringent heat inputs levels in order to maintain the acceptable impact energy levels. Furthermore, since welding productivity levels are associated with heat input, the cost effectiveness of HSLA-65 steels may be offset by these stringent heat input levels. Future plans: Further fracture toughness testing is planned to investigate the influence of the notch placement, along with additional welding trials using novel higher productivity welding processes.

DRDC Atlantic TM 2008-130 iii

Sommaire

The Influence of Heat Input on the Fracture and Metallurgical Properties of HSLA-65 Steel Welds: Fabrication conditions, Tensile, Impact and Microstructures

Christopher Bayley; Adam Mantei; DRDC Atlantic TM 2008-130; R & D pour la défense Canada – Atlantique; Octobre 2008.

Introduction : L’acier du type HSLA-65 (norme ASTM A945, nuance 65) constitue un acier présentant une très bonne soudabilité, caractérisé par sa faible teneur en carbone et en éléments d’alliage, et conséquemment, sa faible teneur en équivalent carbone. L’acier de ce type a été élaboré afin d’occuper un créneau se situant entre les aciers extrêmement résistants du type HSLA-80, utilisés pour fabriquer des canalisations, et les aciers classiques, très résistants, qui sont employés pour exécuter les travaux sur les chantiers maritimes (c.-à-d. ceux du type DH/EH-36, selon la certification de l’American Bureau of Shipping). L’élaboration de l’acier HSLA-65 visait à obtenir un alliage présentant une résistance mécanique, une résilience et une soudabilité supérieures, et pouvant servir à fabriquer des structures de navires militaires. Les propriétés des aciers de ce type dépendent de leur faible teneur en carbone et de la nature des éléments de microalliage, ainsi que des conditions de traitement thermomécanique, plutôt que des conditions propres aux procédés d’alliage comportant des étapes de trempe ou de revenu, qui sont assez coûteux. Les résultats d’études des structures de coques de bâtiments de la classe des destroyers indiquent que les avantages que présentent des aciers de 80 ksi (milliers de lb/po2), en matière de résistance, sont limités par le phénomène de flambage des plaques, et qu’un acier caractérisé par une limite élastique proportionnelle de 65 ksi peut résister à une charge semblable. De plus, les coûts des matériaux et ceux de fabrication de structures à base d’acier HSLA-65 sont comparables à ceux associés aux aciers très résistants employés dans les chantiers maritimes.

Résultats : Le présent rapport contient un résumé des conditions de soudage dans lesquelles ont été fabriqués cinq panneaux soudés distincts, en utilisant le procédé de soudage à l’arc sous protection gazeuse avec fil plein (procédé GMAW) ou celui de soudage à l’arc avec fil fourré (procédé FCAW), ainsi que de leurs propriétés de résistance à la traction, de dureté et de résistance aux chocs. Les travaux comprennent l’étude des panneaux obtenus avec le procédé FCAW et la détermination des effets de l’apport de chaleur lors du soudage et de la nature de l’électrode consommable, ainsi que la comparaison de ces résultats avec ceux des joints soudés obtenus avec le procédé GMAW, pour des valeurs intermédiaires en matière de résistance de l’électrode et d’apport de chaleur.

Les valeurs de résistance à la traction et de dureté de tous les panneaux soudés sont satisfaisantes, mais on a observé un important déplacement de la température de transition ductile-fragile dans le cas de la zone thermiquement affectée (ZTA) des soudures produites avec un apport de chaleur élevé. Il a été établi que cette modification remarquable du comportement de transition est fonction de la propagation de la fissure au sein du gradient de microstructure de la ZTA. La propagation de la fissure dans des régions présentant une microstructure différente peut aussi expliquer la dispersion accrue des données associées au comportement de transition ductile-fragile des soudures produites avec un apport de chaleur élevé.

iv DRDC Atlantic TM 2008-130

Portée : Les résultats fournis dans le présent rapport soulignent le besoin de réguler rigoureusement l’apport de chaleur lors du soudage afin d’obtenir des soudures qui présentent une résistance aux chocs acceptable. De plus, comme le degré d’efficacité du soudage est lié à l’apport de chaleur, le rapport coût-efficacité de l’utilisation des aciers du type HSLA-65 pourrait être compensé par les effets des mesures visant à appliquer ces critères rigoureux à l’apport de chaleur. Recherches futures : On prévoit exécuter d’autres essais de résistance à la rupture afin d’étudier les effets de la position de l’entaille, ainsi que des essais de soudage additionnels, en utilisant des procédés de soudage de pointe offrant un degré d’efficacité supérieur.

DRDC Atlantic TM 2008-130 v


vi DRDC Atlantic TM 2008-130

Table of contents

................................................................................................................................. i Abstract ……...... ........................................................................................................................................ i Résumé

........................................................................................................................ iii Executive summary. ...................................................................................................................................... iv Sommaire

........................................................................................................................... vii Table of contents................................................................................................................................. ix List of figures

.................................................................................................................................. xii List of tables ............................................................................................................................... 1 1 Introduction ........................................................................................................... 4 2 Welding and Fabrication

2.1 .................................................................................................................... 4 FCA Welds2.1.1 ............................................................................................... 5 Plate Inspection

2.2 .................................................................................................................. 6 GMA Welds2.2.1 ........................................................................................ 7 Thermal Monitoring

2.2.1.1 ............................................................................... 7 GMAW Plates2.2.1.2 ................................................................................ 8 FCAW Plates

2.2.2 .................................................................................. 11 Calculated Heat Inputs ............................................................................................ 12 3 Results - Testing and Evaluation

3.1 ..................................................................................... 12 Weld Thermal Time Histories3.2 ................................................................. 15 Metallography and Microhardness Profiles3.3 .......................................................................................... 22 Transverse Tensile Testing3.4 ........................................................................................ 24 Weld Metal Tensile Testing3.5 ................................................................................................ 27 Charpy Impact Testing3.6 ..................................................................................... 30 Prior Austenite Grain Growth

............................................................................................................................... 33 4 Discussion ............................................................................................................................. 39 5 Conclusions

.. ...................................................................................................... 41 Annex A FCAW Plate Usage .............................................................................. 41 A.1 71T1HYN-Low Heat Input Panels ............................................................................. 43 A.2 71T1HYN-High Heat Input Panels .................................................................................. 45 A.3 101TM -Low Heat Input Panels ................................................................................. 47 A.4 101TM –High Heat Input Panels

.. .................................................................................................... 51 Annex B GMAW Plate Usage

.. ........................................................................................... 55 Annex C Weld Thermal Monitoring ............................................................................................................ 55 C.1 GMAW Plate 2 ......................................................................................................... 58 C.2 GMAW – Plate 3

............................................................................................................................... 62 References .....

DRDC Atlantic TM 2008-130 vii

.............................................................................................................................. 63 Distribution list

viii DRDC Atlantic TM 2008-130

List of figures

.......................................................................................................... 5 Figure 1 FCAW Joint Design

Figure 2 Large slag inclusion associated with a repair pass associated with the panel 101TM-HH. ................................................................................................................................ 5

........................................................................................................ 6 Figure 3 GMAW Joint Design

.............................................................. 7 Figure 4 Drill holes for the GMAW plate thermocouples.

Figure 5 Drilled thermocouples on the surface of the GMA plates prior to welding. Thermocouples were located 1.0, 2.5 and 5.0 mm away from the edge preparation. ... 8

............. 8 Figure 6 Outline of GMAW welds and approximate thermocouple locations for Plate 2.

Figure 7 FCAW vertical thermocouple holes drilled from back side and terminating at plate mid-plane....................................................................................................................... 9

Figure 8 Thermocouple positions: a) HYN-HH; b) HYN-LH; c) 101TM HH; d) 101TM LH. The labels “tc” denotes the channel used to acquire data for the thermocouple on channel 3. .................................................................................................................... 10

.................................................................. 13 Figure 9 Weld Plunge Thermal History (101TM-HH)

Figure 10 Thermal profiles from the first pass of 101TM HH. The T8-5 is the time required to cool between 800 to 500 ....................................... 14 oC and summarized in Table 5.

........................................................................................ 14 Figure 11 Third pass - GMAW-Plate 2.

Figure 12 Microhardness traverse across the mid-plane of the GMAW plat e 3. Photographs of the microstructural regions identified as a, b, and c are shown in Figure 13.......... 16

Figure 13 Microstructural regions identified by the labels a, b, c and d in Figure 12. The diamond shaped object in the bottom of each photo are micro-hardness indents. ...... 18

............................ 20 Figure 14 Microhardness (a) and micrograh ( b) of the ER70S-6 weld metal.

Figure 15 Microhardness comparison for the 5 weld metals. All of the traverses start on the bevel edge preparation and terminate on the straight fronted edge preparation.......... 21

Figure 16 HAZ hardness and width of the straight fronted side. The final point represents the extent of the visible HAZ. ........................................................................................... 21

Figure 17 Strain distributions at yield and peak load during a transverse tensile test of the weld. The weld is located in the center of the specimen............................................. 23

............................................... 25 Figure 18 Room temperature engineering stress-strain behaviour.

Figure 19 Macrographs of a) HYN-LH and b) 101TM-LH showing the region sampled by the reduced cross sectional area of the all weld tensile specimen. The relative area fractions of the passes are listed in Table 6................................................................. 25

............................................................ 27 Figure 20 Charpy notch positioning (sketch not to scale).

........................................... 28 Figure 21 HAZ Charpy Impact Energy for the full sized specimens.

DRDC Atlantic TM 2008-130 ix

...................................................................................... 29 Figure 22 Weld Charpy impact energies.

............... 29 Figure 23 Weld and HAZ Charpy Impact Energy from subsize (10x7.5 mm) samples.

Figure 24 Average and 95% confidence intervals of prior-austenite grains diameters measured in the heat affected zone. Data points offset for clarity. ............................ 31

Figure 25 Microstructures sampled by Charpy Impact Notches for the high heat input (a) and low heat inputs welds (b) ............................................................................................ 32

Figure 26 Charpy Impact energy for HYN-HH notched 1.0 mm from the fusion line. Red markers indicated samples which were selected for sectioning while the labelled specimens are reported herein and referred to as the upper and lower shelf specimens. ................................................................................................................... 34

Figure 27 Fracture surface of specimen displaying upper-shelf behaviour as highlighted in Figure 26 (160J @-5oC and 60% cleavage appearance) b) transverse cross section along the centerline. Arrows in a) and b) point to secondary cracks shown in Figure 29 ..................................................................................................................... 36

Figure 28 Fracture surface of specimen displaying lower-shelf behaviour as highlighted in Figure 26 (60J @0oC and 90% cleavage appearance) b) transverse cross section along the centerline. .................................................................................................... 37

.................................. 38 Figure 29 Micrograph of secondary cracks in the upper bainitic structure.

Figure 30 As received panel showing the relative location of the three blocks labelled A,B and C. .......................................................................................................................... 41

........................ 42 Figure 31 Specimen locations and labels from Block A of panel 71T1HYN-LH.

......................... 42 Figure 32 Specimen locations and labels from block B of panel 71T1HYN-LH.

............................ 43 Figure 33 Sample locations and labels from Block C of panel 71T1HYN-LH.


....................... 44 Figure 35 Specimen locations and labels from Block A for panel 71T1HYN-HH

....................... 44 Figure 36 Specimen locations and labels from Block B for panel 71T1HYN-HH

....................... 45 Figure 37 Specimen locations and labels from Block C for panel 71T1HYN-HH


............................ 46 Figure 39 Specimen locations and labels from Block A for panel 101TM-LH

............................. 46 Figure 40 Specimen locations and labels from Block B for panel 101TM-LH

............................. 47 Figure 41 Specimen locations and labels from Block C for panel 101TM-LH


............................ 48 Figure 43 Specimen locations and labels from Block A for panel 101TM-HH

............................ 48 Figure 44 Specimen locations and labels from Block B for panel 101TM-HH

x DRDC Atlantic TM 2008-130

............................ 49 Figure 45 Specimen locations and labels from Block C for panel 101TM-HH

Figure 46 GMAW Plate 1. The central region contained the drilled holes for the thermocouples. ............................................................................................................ 51

........................................ 52 Figure 47 GMAW Plate 2 showing the location of various specimens.

....................................... 53 Figure 48 GMAW Panel 3 showing the location of various specimens.

.......................................................................................... 55 Figure 49 First Pass – GMAW Plate 1

....................................................................................... 56 Figure 50 Second Pass- GMAW Plate 2

......................................................................................... 57 Figure 51 Third Pass – GMAW Plate 2

.......................................................................................... 58 Figure 52 First Pass – GMAW Plate 3

...................................................................................... 59 Figure 53 Second Pass – GMAW Plate 3

......................................................................................... 60 Figure 54 Third Pass – GMAW Plate 3

....................................................................................... 61 Figure 55 Fourth Pass – GMAW Plate 3

DRDC Atlantic TM 2008-130 xi

List of tables

Table 1 Chemical analyses (wt%). Single values in the first row represent specified maximum. Blank values indicate that the composition was below the detection limit. .............................................................................................................................. 2

.............................................................................. 4 Table 2 FCAW Panel Labels and Designations

Table 3 Location of thermocouples acquired from transverse sections at the thermocouple location which yielded valid data................................................................................ 11

................................................................. 11 Table 4 Calculated Heat Inputs (kJ/mm) for each pass

Table 5 Recorded Cooling times from 800-500oC. FL= Fusion Line, WM=weld metal plunges. ....................................................................................................................... 15

Table 6 Relative area fractions (%) of weld passes sampled by the all weld tensile specimen..... 24 ................................................................................. 26 Table 7 Material properties – tensile testing

Table 8 HAZ Ductile-Brittle Transition Temperatures assessed using the 50% FATT and mid-shelf impact energy.............................................................................................. 30

Table 9 Weld Metal Ductile-Brittle Transition Temperatures assessed using the 50% FATT and mid-shelf impact energy. ...................................................................................... 30

xii DRDC Atlantic TM 2008-130

1 Introduction

HSLA-65 (ASTM A945 Grade 65) is a highly weldable steel that is characterized by its low carbon and alloy content, and hence carbon equivalent. It was designed to fill the niche between higher strength HSLA-80 pipeline steels and conventional shipyard high-strength steels (i.e. American Bureau of Shipping DH/EH-36). HSLA65 was developed to have improved strength, toughness and weldability for naval ship structures. These steels depend on low carbon content and microalloying additions, in conjunction with thermo-mechanical processing, to obtain their properties, rather than relatively expensive alloying and/or quench and tempering. Design studies of destroyer-type hulls showed that the strength benefits of 80 ksi steels were limited by plate buckling and that a 65 ksi yield strength steel could achieve similar weight. Furthermore, fabrication and material costs of HSLA-65 are similar to shipyard high-strength steels. The replacement of HSS with HSLA-65 allows thinner structures to be created, and hence smaller volumes of weld material, thus reducing the fabrication costs.

Following the alloy’s introduction, a procurement specification (ASTM A945[1]) was developed which allows plate from a range of manufacturing processes for plates up to 32 mm (1 1/4”) thick. The specified chemical composition is listed in Table 1, along with the plate product analysis and the chemical composition of the welding electrodes. The chemical requirements are based on the line-pipe grade X70 [2] and intrinsically related to the effects of processing conditions on microstructure development and weldability For example, the carbon content is limited to 0.1 wt% to ensure weldability, manganese content ranges from 1.1 to 1.65 wt% to provide adequate toughness and strength, silicon is added for deoxidation and solid solution strengthening. The maximum limits for nickel, chromium, copper and molybdenum are imposed to prevent excessive increases in strength and hardenability and carbon equivalent. Microalloying additions of columbium, vanadium and titanium form a solid solution in austenite and increase the austenite recrystallization temperature, thereby limiting austenite grain growth. The specified range for titanium fixes solute nitrogen as TiN rather than TiC which is associated with brittle behaviour. The presence of TiN precipitates limits austenite grain growth during welding, preserving heat affected zone (HAZ) strength and toughness. Optimum grain boundary pinning is obtained with TiN particles less than 0.05 μm with titanium content of about 0.015% with a stoichiometric ratio of Ti to N of no more than 3.5 [3, 4]. Aluminum serves both as a deoxidant, grain refiner and fixant for solute nitrogen, which minimizes strain ageing effects [5].

DRDC Atlantic TM 2008-130 1

Table 1 Chemical analyses (wt%). Single values in the first row represent specified maximum. Blank values indicate that the composition was below the detection limit.

C Mn Si Ni Cr Cu Mo Ti Al N S ASTM A945

0.1 1.1 -1.65

0.1 -0.4

0.4 0.2 0.35 0.08 0.007-0.020

0.08 0.012 0.01

5/8” Plate 0.08 1.44 0.308 0.02 0.04 0.019 0.01 0.015 0.02 0.006 0.004 3/8” Plate 0.08 1.43 0.299 0.01 0.04 0.020 0.006 0.016 0.03 0.008 0.004

As Deposited 71HYN-

LH

0.09 1.5 0.3 0.34 0.05 0.03 0.006 0.16 0.007

As Deposited

0.09 1.5 0.3 0.25 0.05 0.03 0.007 0.14 0.007

71HYN-HH As

Deposited 101TM-

LH

0.09 1.4 0.28 1.4 0.04 0.03 0.006 0.17 0.009

As Deposited 101TM-

HH

0.08 1.2 0.24 1.3 0.04 0.03 0.005 0.11 0.009

As Deposited ER70S6

0.09 1.4 0.95 0.034 0.022 0.006 0.013

HSLA-65 is considered to have excellent weldability characteristics. Steel weldability is a function of the hardenability of the steel, which can be related to the steel composition and, in particular, to the carbon content. Many expressions have been developed to relate the hardenability to the material chemistry, with the two most common being the IIW carbon equivalent (CEIIW) and the Pcm index. The Pcm index was developed for pipeline steels with carbon contents less than 0.11% (weight percent), which is outside of the range of the CEIIW formulation [6].

1556CuNiVMoCrMnCCEIIW

++

++++= Eq 1)

BVMoNiCuCrMnSiCPcm 51015602030

++++++

++= Eq 2)

Regardless of the hardenability expression used, HSLA-65 can welded without additional pre-heat requirements. However, the hardenability index does not address the susceptibility of the base material to the weld thermal cycle. The objective of this welding study was to determine the influence of welding heat input on the mechanical properties of the weld metal and heat affected

2 DRDC Atlantic TM 2008-130

zones. Mechanical testing included all weld metal tensile, transverse tensile and Microhardness, along with the determination of the Charpy Impact transition behaviour. For this assessment, weld panels were fabricated using either the flux core are welding (FCAW) process or gas metal arc welding (GMAW). For the FCAW process, two different consumables and two different welding heat inputs were used. The base material for the following study was a controlled rolled 15.8 mm (5/8”) thick plate for the FCAW welds and a 9.375 mm (3/8”) thick plate for the GMAW welds. In all cases, the weld line ran transverse to the plate rolling direction.


2 Welding and Fabrication

The specimen set consisted of two welding processes: Flux-Core Arc Welding (FCAW) and Gas Metal Arc Welding (GMAW) with varying heat inputs and weld consumables. The FCAW panels were welded under contract W7707-063485/001/HAL, while the GMAW welds were fabricated at the Fleet Maintenance Facility CAPE BRETON (FMF CB) welding shop.

In order to determine the cooling rates at various locations from the fusion line, K-type thermocouples were staggered along the plate mid-thickness by drilling from either the back or front face of the plates. The thermocouples were positioned at various distances away from the initial edge preparation and the time-temperature traces were recorded for each pass.

2.1 FCA Welds

For the FCAW specimens, a matrix of four different welding configurations was used, involving variations in the weld heat input and the electrode classification:

Table 2 FCAW Panel Labels and Designations

Label Electrode Heat Input HYN-LH MIL-71T1-HYN Low Heat HYN-HH MIL-71T1-HYN High Heat 101TM-LH MIL-101TM Low-Heat 101TM-HH MIL-101TM High-Heat

MIL-71T1-HYN and MIL-101TM are the two filler metal classifications according to MIL-E-24403/1D and MIL-E-24403/2B, respectively. These flux core electrode designations state that the minimum tensile strength of the weld metal is either 70 or 100 ksi and that the electrode can be used in all positions with a 75% argon 25% CO2 shielding gas. The HYN designation indicates a low hydrogen (5.0 mL per 100 g of weld) content.

High heat input, multi-pass welds correspond with larger weld beads and hence fewer welding passes. For the high heat input welds, 4 passes were required to fill the joint, while the low heat input welds required 8 passes. The calculated heat inputs for each pass are listed in Table 4.

The panels were welded in the vertical up direction. To ensure a constant heat input along the entire weld length of 1828 mm (72”), a mechanized gantry was employed. Prior to welding the bevel joint, the panels were affixed with strong backs in order to restrict the amount of weld induced distortion.

For the FCAW panels, butt-welds were made between plates of thickness 5/8” (15.88mm). The joint configuration was a single bevel configuration with a root opening of 5.0 mm as illustrated in Figure 1. This joint preparation was chosen in order to promote a straight-fronted HAZ in order to facilitate sampling a single microstructural region in the notch of the Charpy Impact specimen.


Sketches of the as-received panels are shown in Appendix A. These sketches include the measured location of the thermocouples from the edge of the root pass along the location and the label of each test specimen.

45.00°5.00 mm

15.88 mm Base BaseWeld

Figure 1 FCAW Joint Design

2.1.1 Plate Inspection

Upon receipt of the plates at Dockyard Laboratory Pacific, they were radiographed and visually inspected. A number of defects were found that were associated with weld repair, slag entrapments or thermocouple plunges. The plunged thermocouple introduced foreign material into the weld pool generating porosities and other local discontinuities. Furthermore, the plunges created a local disruption to the slag pool which caused the molten weld metal to spill out in a manner that was exacerbated by the vertical-up welding position [7].

For the HYN-HH panel a repair of some incomplete penetration resulted in a weld repair, which was noted on the surface of the panel, and periodic discontinuity likely associated with a slag entrapment. A large nearly continuous slag entrapment associated with a weld root repair was found on the 101TM-HH panels as shown in Figure 2. However, since the defect was located close to the bottom surface, it would be partially removed during the subsequent machining operations used to extract test specimens. Nevertheless, remnants of the slag inclusion were still visible on some of the test coupons, particularly the Drop Tower specimens. In these cases, the notch was located on the same side as the inclusion in order to minimize its influence on the fracture process.

Figure 2 Large slag inclusion associated with a repair pass associated with the panel 101TM-

HH.


2.2 GMA Welds

In parallel to the welding of the FCAW plates, FMF CB was tasked to fabricate four 508 mm welded panels using the GMAW process. The plates used in the GMAW testing program were 3/8” (9.5mm) thick, with an HSLA-65 backing plate to support the root welding pass, as illustrated in Figure 2. The welds were fabricated in the flat 1G position and used an AWS ER70S-6 together with a 75% Argon 25% CO gas mixture. 2

Sketches of the as-received plates are included in Appendix B and include the location and label of each test specimen which was machined from the panels.

9.26 mm

60.34°

Base Base

Weld

Figure 3 GMAW Joint Design


2.2.1 Thermal Monitoring

In order to track the thermal cycle experienced by various locations in the base plate, thermocouples were placed away from fusion line and located at the plate’s mid-plane. Thermocouples were also plunged into the still-molten weld metal to measure its cooling rate.

2.2.1.1 GMAW Plates

For the GMAW plates, thermocouple holes were drilled into the plates at a 30° angle, and terminated 1.0, 2.5 and 5.0 mm away from the edge as shown in Figure 4. Into these holes K-type thermocouples covered with 3/8” ceramic insulation were spot welded to the base of each hole. A 4th reference thermocouple was also spot welded onto the surface of the plate, about 5.0 mm away from the edge. These four thermocouples are shown in 6Figure 5, while Figure is a slice through Plate Two at the 5.0 mm thermocouple which shows the approximate location of the other two thermocouples and the weld pass profiles.

The temperature-time history was recorded using a Pace Scientific XR5-SE-M-50mV Data Logger at a rate of 25 Hz. At this collection rate, the data is stored on internal memory and down loaded to a PC between the welding of each plate, thus eliminating the possibility of viewing the thermal data in real-time. Due to a malfunctioning trigger level used to initiate data collection, the thermal data for the first panel was not acquired.

31.00°118.00°

0-3/16"

1.000 mm

0-1/8"0-3/8"

0-3/8"

Figure 4 Drill holes for the GMAW plate thermocouples.


Figure 5 Drilled thermocouples on the surface of the GMA plates prior to welding. Thermocouples were located 1.0, 2.5 and 5.0 mm away from the edge preparation.

Figure 6 Outline of GMAW welds and approximate thermocouple locations for Plate 2.

2.2.1.2 FCAW Plates

The locations of the thermocouples were staggered and positioned to nominally collect the thermal histories at three different locations: the fusion line (FL), 1.5 mm into the heat-affected


zone (HAZ), and 4 mm into the HAZ. Unlike the GMAW thermocouple holes, the FCAW thermocouple holes were drilled from the backside and terminated at the plate mid-thickness as sketched in Figure 7. In addition thermocouples were plunged into the molten weld metal in order to record the weld thermal profile.

However, due to the unforeseen extent of weld dilution, whereby the molten region of the weld extends into the base plate, many thermocouples did not sample their anticipated location. Figure 8 shows the representative locations of the thermocouples with relation to the weld fusion line for the FCA panels. The locations were obtained from transverse cross sections through selected thermocouples, however not all the thermocouple locations could be obtained, because some thermocouple holes had been cut from the plate and mounted for microscopy without retaining information on exactly which thermocouple it represented. A summary of the thermocouple locations which yielded valid thermal data is shown in Table 3.

For the FCAW panels, the thermocouple wires were twisted the entire length of the hole, instead of being inserted straight. Since the thermocouples measure the change in resistance between two metallic wires, the recorded measurements are averages obtained from each contact point along the length of twisted pair.

1.5 to 4.0 mm

15.8

0 m

m

2.30 mm Figure 7 FCAW vertical thermocouple holes drilled from back side and terminating at plate mid-

plane.


a) b)

c) d)

Figure 8 Thermocouple positions: a) HYN-HH; b) HYN-LH; c) 101TM HH; d) 101TM LH. The labels “tc” denotes the channel used to acquire data for the thermocouple on channel 3.


Table 3 Location of thermocouples acquired from transverse sections at the thermocouple location which yielded valid data.

Expected Location (tip to FL on midplane)

(mm)

Distance from Final Fusion

Line Closest

Pass Test Panel tc (mm)

0 0 1 9 HYN HH 1.5 0 1 10

0 0 1, 2 1 HYN LH 1.5 0 2 2

0 0 1 1 1.5 0 1 2

9 0 0 1 101TM HH

1.5 0.5 1 10 0 0 2 1

1.5 0.5 2 2 9 0 0 2

101TM LH

1.5 0 2 10

2.2.2 Calculated Heat Inputs

The calculated heat inputs for all of the panels and processes are summarized in Table 4. The heat input (Q) was determined from the arc voltage (A) and current (V) as follows:

LtVAmmkJQ⋅⋅⋅

=1000

)/(

where t is the time in seconds and L the distance travelled. The heat input is used to determine the amount of energy introduced into the material and is directly related to the HAZ thermal cycle and subsequent microstructural changes.

Table 4 Calculated Heat Inputs (kJ/mm) for each pass

Welding Pass Process Consumable Label 1 2 3 4 5 6 7 8Mil 71T1-HYN HYN-LH 1.85 1.46 0.91 1.02 1.02 1.05 1.05 1.05Mil 71T1-HYN HYN-HH 4.24 1.85 1.41 1.36

FCAW

Mil 101TM 101TM-LH 1.82 1.34 0.91 1.02 1.22 1.05 1.05 1.05Mil 101TM 101TM-HH 4.32 1.97 1.41 1.43

GMAW AWS ER70S-6 Panel 1 2 1.8 0.75 1.25GMAW AWS ER70S-6 Panel 2 NA 1.8 NA NAGMAW AWS ER70S-6 Panel 3 NA NA NA NA


3 Results - Testing and Evaluation

In order to assess the structure-property relationship of the HAZ material, metallurgical evaluations and mechanical testing were performed. The metallurgical assessment included optical microscopy and micro-hardness traverses, while the mechanical testing involved tensile strength and impact energy measurements.

3.1 Weld Thermal Time Histories

Sample temperature-time history plots for plunged thermocouples and drilled thermocouples for the 101TM-HH plate are shown in Figure 9 and Figure 10, respectively, while a sample thermocouple time history for the third pass of Plate 2 GMAW is plotted in Figure 11.

The plunged thermocouple temperature-time histories typically resulted in smooth and consistent temperature profiles. This consistency is especially evident in Figure 9 for the two plunged thermocouples acquired from the second pass. The plunged thermocouples provide an efficient approach to measuring the cooling rate, but do not capture the full thermal cycle. In particular, the plunged thermocouples fail to capture the rising and dwell time during which thermally dependent processes, such as precipitate ripening and dissolution along with grain coarsening, can occur. This thermal data can be inferred from the drilled thermocouples.

The thermal profiles along the fusion line (FL) and at various points within the HAZ for the first pass of the 101TM-HH panel are plotted in Figure 10 and for the GMAW weld in Figure 11. These thermal histories are influenced by their proximity to the fusion line, which influences both the peak temperature and the heating and cooling rates. For both panels, the heating and cooling rates are a function of the speed of the moving heat source and the panel geometry. For all of the measurement locations, the temperature is influenced by the approaching and retreating heat source, with significant heating occurring well ahead and behind the torch location. This heating ahead of the torch accounts for the significant times required to achieve the peak temperature. For the case of the FCAW panels, the heating rate ahead of the torch is 5-10 time faster than for the GMAW panels. Differences in arc efficiency, molten metal transfer and welding position (vertical-up versus flat for the FCAW and GMAW, respectively) all contribute to this noted difference. Similarly, the cooling rates are also process dependent, with most of the significant difference being the presence of an insulating flux covering instead of the forced convection associated with the gas shield environment around a GMAW torch.

However, some of the differences between the two thermal profiles are believed to be setup dependent. This setup dependence originates from the presence of different sized drill holes, thermocouple placement within a highly non-linear thermal gradient and different thermocouple junctions (i.e. twisted versus straight). The latter is particularly evident for the FCAW panels, whereby the temperature would have been averaged along the length of the twisted junction. This implies that the peak temperatures for the first pass were always higher than subsequent passes, regardless of the distance from the thermocouple tip to the weld boundary. However, despite the averaging influence of the twisted thermocouple junctions, the agreement between the measured cooling times from 800-500C for the plunged and drilled thermocouple were surprisingly good. Apart from the first pass of 101TM-HH, the difference between the weld plunge and drilled


fusion line thermocouples is only 2-3 seconds. This excellent agreement may be in part due to the diminishing thermal gradient as the torch retreats from the measurement location. The cooling times between 800 and 500 C are particularly important since this time defines the final transformation product from the austenite grains present at high-temperatures. These cooling times for all the panels are given in Table 5.

Regardless of the thermocouple technique, the cooling time between 800-500 C was noticeably slower in the high heat input cases (50-70 seconds) than the cooling times in the low-heat input cases (10-20 seconds).

0

200

400

600

800

1000

1200

1400

0 20 40 60 80 1Time (s)

Tem

pera

ture

(C)

00

Ch5 Pass 1 (T8-5=36.2 s)Ch4 Pass 2 (T8-5=18.4 s)Ch5 Pass 2 (T8-5=18.3 s)

Figure 9 Weld Plunge Thermal History (101TM-HH)


0

200

400

600

800

1000

1200

0 10 20 30 40 50 60 70 80 90 10Time (s)

Tem

pera

ture

(C)

0

FL (tc1) (T8-5=50.5 s)FL (tc2) (T8-5=55.3 s)HAZ (tc3) (T8-5=NA)HAZ (tc11) (T8-5=NA)

Figure 10 Thermal profiles from the first pass of 101TM HH. The T8-5 is the time required to

cool between 800 to 500 oC and summarized in Table 5.

0

100

200

300

400

500

600

700

800

900

1000

0 10 20 30 40 50 6Time (s)

Tem

pera

ture

(C)

0

T1 (T8-5=6.2s)

T2

T3

Figure 11 Third pass - GMAW-Plate 2.


Table 5 Recorded Cooling times from 800-500oC. FL= Fusion Line, WM=weld metal plunges.

∆T 800-500 time (seconds) Heat

Input FL FL FL FL WM WM (tc1) (t2) (tc9) (tc10) (tc4) (tc5) (kJ/mm)

Pass 1 1.85 62.45 67.9 59 64.8 HYNHH Pass 2 1.46 13.75 13.05 Pass 1 4.24 24.35 21.8 20.45 24.5 24.9 Pass 2 1.85 14.5 Pass 3 1.41 5.2

HYNLH

Pass 4 1.36 8.3 5.2 Pass 1 4.32 50.5 55.3 36.2 101TMHH Pass 2 1.97 18.4 18.3 Pass 1 1.82 13.65 16.85 14.55 Pass 2 1.34 12.7 11.85 Pass 3 0.91 8.15 10.15

101TMLH

Pass 4 1.02 7.2 GMAW P2 Pass 3 6.2 GMAW P3 Pass 3 7.6

3.2 Metallography and Microhardness Profiles

Microhardness characterization of the welds was performed in order to quantitatively assess the influence of the welding thermal cycle on the microstructure. Microhardness traverses across the welds and HAZ were acquired from mounted and polished sections. For each sample, a linear series of measurements was taken every 200 μm in two configurations: the first was approximately parallel to the angled bevel side of the weld, while the second was a series of horizontal traverses at the 1/3, 1/2 and 2/3 thickness positions which extended from the fusion line into the base material.

A sample microhardness traverse across the GMAW P3 weld is plotted in Figure 12 with accompanying microstructures of selected regions in Figure 13. The traverse clearly delineates the difference in hardness associated with the base metal, HAZ and weld metal that was related to the peak temperature and cooling associated with the weld thermal cycle. Traversing from left to right, the base metal has an average hardness of 220 Hv and is comprised of a banded ferritic and pearlitic microstructure as shown in Figure 13a. The banding is a consequence of the controlled rolling schedule which specifies significant plate reduction in the dual phase austenite-ferrite region. Upon entering the subcritical heat affected zone (SCHAZ), there is a drop in the hardness associated with the tempering and stress relieving effect of the peak temperature. Point b in Figure 12 is within the inter-critical heat affected zone (ICHAZ) bounded by the Ac1` and Ac3 temperatures. These temperatures define the dual phase region within which austenite, ferrite and pearlite are present. The Ac1 temperature is the temperature at which ferrite and pearlite begin the phase transformation into austenite, while the Ac3 defines the high temperature limit of ferrite and was estimated as 769 oC [5] for this steel. Above this temperature, the microstructure would be


fully austenite. The microstructure at point b is shown in Figure 13 and is a mixture of fine ferrite and pearlite. The phase transformation from a banded ferrite and pearlite to austenite during the heating portion of the thermal cycle, followed by the phase transformation back to ferrite and pearlite is responsible for the breaking up of this banded structure.

Between points b and c in Figure 12, there is a rapid change in the hardness which corresponds to the coarse grained heat affected zone (CGHAZ) associated with peak temperature above Ac3. As shown in Figure 13c, the microstructure is bainite which formed during the rapid cooling cycle and extends the width of the CGHAZ between the intercritical region up to the fusion line. Bainite is a metastable microstructure resulting from the transformation from austenite at temperatures between those which will form martensite and pearlite. The structure of Figure 13c and d is typical of upper bainite, which consists of plates of cementite in a ferrite matrix. These discontinuous carbide lathes tend to have parallel orientations in the longer direction of the bainite area [8].

190

200

210

220

230

240

250

260

270

-10 -8 -6 -4 -2 0 2 4 6 8 1

Distance from Weld Centerline (mm)

Vick

ers

Har

dnes

s (H

V)

0

Weld metal(Untempered 2nd Pass)

Weld Metal(Tempered 1st Pass)

HAZ BaseHAZBase

a

c

d

b

Figure 12 Microhardness traverse across the mid-plane of the GMAW plat e 3. Photographs of

the microstructural regions identified as a, b, and c are shown in Figure 13.

Beyond the fusion line at point d in Figure 12 and extending for 10 mm is the solidified weld metal. The weld metal is consistently harder than the parent material, but with greater variability. The as-solidified weld metal forms acicular ferrite surrounded by grain boundary ferrite which forms on the prior-austenite boundaries, as shown in the right of Figure 13d. Acicular ferrite has a basket weave morphology and nucleates from a fine dispersion of inclusions within the austenite grain.


a) Banded ferrite microstructure. The black regions are pearlite, the white ferrite.

b) Fine ferritic microstructure within the intercritical region where the peak temperature was between Ac and Ac . 1 3


c) Bainitic microstructure within the coarse grained heat affected zone (CGHAZ). Bainite is composed of cementite laths in a ferrite matrix. In this region the peak temperature was sufficient to form austenite (i.e. above the Ac ) 3

d) Fusion boundary. To the left of the boundary is the coarse grained heat affected zone composed of upper bainite and polygonal ferrite. In contrast, the weld metal to the right of the fusion boundary is mainly acicular ferrite with proeutectoid ferrite (grain boundary ferrite) outlining the prior austenite boundaries.

Figure 13 Microstructural regions identified by the labels a, b, c and d in Figure 12. The diamond shaped object in the bottom of each photo are micro-hardness indents.

The influence of the tempering of the subsequent passes is examined in Figure 14. Figure 14a plots the microhardness values traversing from the top of the 4th pass through 2nd pass and


terminating in the 1st pass. The weld metal which is tempered as a result of the subsequent passes is clearly seen in both the hardness traverse in Figure 14a and the associated micrograph in Figure 14b. In the identified tempering region, the columnar feature of the as-deposited weld metal is broken up, resulting in more equiaxed shaped grains. In the tempered regions, there is a significant drop in the hardness compared with the untempered weld metal.

The traverse microhardness profiles at relative locations in the base, HAZ and weld metal are plotted in Figure 15 for the 5 welded test panels. In all cases, the hardness profiles cross the midpoint of the first pass. In all of the panels, the weld metal hardness is greater than the base material. The weld hardness of 101TM-HH is significantly greater than the other welds; macrosegregation of alloying constituents may contribute to the significant increase in hardness along the weld centerline.

Hardness profiles acquired from the averages of three horizontal profiles, which start at the fusion line boundary of the straight fronted side and extent into the HAZ, are plotted in Figure 16 for the 5 welded panels. The relative widths of the heat affected zones are clearly visible and follow the trend of decreased HAZ width with decreasing heat input.


190

200

210

220

230

240

250

260

270

0 2 4 6 8 10 12 14 16Distance (mm)

HV

4th Pass TemperedRegion

2nd Pass 1 st PassTemperedRegion

a) Microhardness traverse extending from the center of the 4th pass through the 2nd and 1st passes. The arrows correspond to the arrows shown in b)

b) Micrograph showing the tempered region between the 4th and 2nd pass. The arrows correspond with indentations above.

Figure 14 Microhardness (a and micrograh ( b) of the ER70S-6 weld metal.


150

175

200

225

250

275

300

Relative Position

Vick

ers

Har

dnes

s

101TM-HH101TM-LHHYN-HHHYN-LHER70S-6

Base HAZ Weld Metal HAZ Base

Figure 15 Microhardness comparison for the 5 weld metals. All of the traverses start on the

bevel edge preparation and terminate on the straight fronted edge preparation.

150

175

200

225

250

275

300

0 0.5 1 1.5 2 2.5 3 3.5

Distance From FL (mm)

Vick

ers

Har

dnes

s (H

V)

4

ER70S-6101TM-HH101TM-LHHYN-HHHYN-LH

Figure 16 HAZ hardness and width of the straight fronted side. The final point represents the

extent of the visible HAZ.


3.3 Transverse Tensile Testing

Transverse tensile tests were performed to determine the relative strengths of the base and parent materials and were evaluated by failure and necking location. Failures occurring in the weld metal result from a weaker weld metal and are associated with strain localization in the weld metal. When weld metal strength is less than the base material, it is labelled “under-matched”, while failures which occur in the base material are labelled “over-matched”. Generally under-matching is undesirable since the fracture toughness of the weld metal must be significantly greater to prevent the onset of crack propagation [9].

One specimen from each welded panel was prepared with the weld positioned perpendicular to the loading direction. During the test the displacement of the actuator, the applied force and a series of images were acquired to determine the relative strain concentration in the weld and base metals. The strains were acquired from a series of 5.0 mm spaced dark lines marked onto the face of each specimen. These faces were subsequently digitally photographed every second throughout the duration of the test. Through a series of image manipulation sequences, the distance between each adjacent line was calculated, thus allowing strain to be determined along the length of the specimen. A representative strain profile acquired at yield and at the peak load of HYN-LH is plotted in Figure 17. In this figure, a color bar represents the true strain measured between adjacent black lines. The weld is located in the central portion of each figure but undergoes little strain compared to the base material and quantifiably demonstrates the over-matching of the welds.

For all of the five welded panels, the weld and adjacent region experienced limited strain in comparison to the base plate. The location of the neck always coincided with the longer length of base metal in the gauge length. This was influenced by both the location of the weld towards the center of the specimen, and the asymmetrical bevel joint design. Upon further examination, there was minimal strain not only within the weld, but in a 10 mm region adjacent to it. The lack of strain adjacent to the weld results from the over matched weld strength, which imposes an additional constraint in the material surrounding the weld. This additional constraint pushes the localization away from the weld and into the base material.


a) Strain at yield b) Strain distribution at Peak Load

c) True strain Colour bar

Figure 17 Strain distributions at yield and peak load during a transverse tensile test of the weld. The weld is located in the center of the specimen


3.4 Weld Metal Tensile Testing

Of special interest was the effect of differing heat inputs and welding electrode classification on the weld metal tensile behaviour. These tensile tests were conducted for each specimen type at least once at room temperature and at approximately -40oC.

A comparison of the room temperature engineering stress-strain data is provided in Figure 18. The plots clearly show the relative differences in strength between the base and the weld metal which was inferred from the transverse tensile tests. Evident as well is the difference in strength for the high and low heat input conditions, particularly for the 101TM electrode cases, while the HYN series electrodes are far less sensitive to the heat input. For the 101TM electrodes the low heat input welds consistently have higher weld metal flow stress than the high heat input panels for the same electrode classifications. However, the relative location of the weld passes sampled in the tensile specimen was different for each panel. As a consequence, the volume fraction of tempered versus un-tempered weld metal in each specimen was different. For example, specimens machined from the HYN-LH panels included a significant portion of the 8th pass which was un-tempered, while the material sampled in the 101TM-LH panels had been tempered by subsequent passes.

Table 6 Relative area fractions (%) of weld passes sampled by the all weld tensile specimen.

Weld Pass 1 2 3 4 5 6 7 8 HYN-LH 1 62 22 15 HYN-HH 24 50 26 101TM-LH 22 31 30 16 101TM-HH 46 53


0

100

200

300

400

500

600

700

800

900

0 0.05 0.1 0.15 0.2 0.25 0.3Engineering Strain

Engi

neer

ing

Stre

ss (M

Pa)

HSLA-65 Base

101TM-HH 71-T1 HYN-HH

101TM-LH 71-T1 HYN-LH

ER70S-6

Figure 18 Room temperature engineering stress-strain behaviour.

a) 71T-1-HYN -LH b) 101TM-LH

Figure 19 Macrographs of a) HYN-LH and b) 101TM-LH showing the region sampled by the reduced cross sectional area of the all weld tensile specimen. The relative area fractions of the

passes are listed in Table 6.

The low temperature tests were conducted in order to asses the temperature dependence of the flow stress. The testing procedure involved cooling the air in the environmental chamber to within ±2 C of -40 C. After holding for a few minutes at this temperature, the extensometer was


attached to the specimen. The time required to open the chamber to insert the extensometer, close the chamber and start the test was approximately 20 seconds. The reason for limiting the exposure of the extensometer to the cold test temperature was to minimize ice build-up on the extensometer. However, as a result of this concern, the stated temperature of -40oC is only approximate, since opening the environmental chamber caused a significant temperature fluctuation.

Table 7 summarizes the engineering stresses at yield and ultimate tensile strength for the room and -40oC temperatures. The results indicate that when cooled to -40 C, the specimens generally exhibited higher ultimate tensile strengths and higher tensile strains at necking, while the yield strengths remained similar, or in some cases, even decreased.

Table 7 Material properties – tensile testing

Young's Modulus

Upper Yield Strength

Lower Yield Strength

Ultimate Tensile Strength Weldment Spec. Temperature

deg C E (GPa) S (MPa) S (MPa) UTS (MPa) YU YL

*No strain data acquired. The weld metal tensile data for ER70S-6 meet the requirements for an ABS grade 4Y400 consumable [10] (Yield > 400 MPa, Tensile Strength 510-690 MPa and elongation > 22%) which would be a suitable consumable used to join HSLA-65 steels. However, the consumables selected for the FCAW process exceed the tensile strength specification for the 4Y400 classification.

BT1 21.3 219.6 703.2 670.7 735.4 BT2 -40 192.1 710.0 671.4 758.1

HYN-HH

-41 193.5 689.0 661.7 770.4 BT3 AT1 21.3 183.2 701.2 701.2 755.7 AT2 21.3 213.0 690.0 676.9 744.6

HYN-LH

-40 190.0 733.0 718.2 806.0 AT3 -42 174.7 752.1 723.7 807.7 AT4

DT1 21.3 237.7 670.7 624.0 694.8 DT2 21.3 186.5 648.1 623.6 694.0

101TM-HH

-40 206.3 708.9 661.9 745.8 DT3 -40 264.2 705.7 651.0 755.1 DT4

CT1 21.3 204.3 745.0 712.3 755.8 CT2 21.3 198.5 720.2 693.0 742.1

101TM-LH

-40 191.5 709.9 692.4 773.3 CT3 -40 180.8 698.1 698.1 808.1 CT4

AWM1 21.3 207.1 605.8 567.9 647.3 AWM2 -40 N/A* 597.9 563.3 669.7

ER70S-6

-40 173.4 564.1 558.8 667.1 AWM3


3.5 Charpy Impact Testing

For each of the 5 different weld panels, two sets of 16 or more Charpy V-Notch specimens were machined. One set of 16 had its notch located 1.0 mm into the HAZ, parallel to the fusion line, while the other had its notch cut right in the centre of the weld. The notch ran normal to the plane of the plates in all cases, with the crack propagating in the longitudinal direction of the weld as shown in Figure 20. Due to the plate thickness used for the GMAW welds subsize (7.5×10 mm) specimens were used rather than the typical 10x10 specimens. Unfortunately, due to the difference in constraints associated with sub-size specimens, their results cannot be directly compared with full scale specimens [11]. Consequently, the results from the GMAW panels are shown separately.

1 mm

HAZ Notch Weld Notch

Figure 20 Charpy notch positioning (sketch not to scale).

The Charpy test provides a qualitative but reliable method to assess the weld properties and, in particular, examine the ductile-brittle transition behaviour. For the Charpy testing, several specimens of each weld were tested in a range of temperatures between -80 C and room temperature in an effort to map out ductile-to-brittle transition behaviour.

The Charpy results demonstrated a distinct shift in transition temperature for the HAZ region of the high heat input welded panels compared with those from the low heat input panels. Not only was the ductile-to-brittle behaviour at much higher temperatures for the high heat input welds, but there was also a substantial increase in the amount of scatter in the transition region. The ductile-to-brittle transition temperature was assessed from both the mid-shelf impact energy and independently from the fracture appearance transition temperature. The mid-shelf impact energy was taken as the temperature associated with the average upper and lower energy shelves, while fracture appearance transition temperature was assessed from the degree of crystallinity on the fracture face in accordance with ASTM E23 [11]. The transition temperatures are summarized in Table 8 for the notched HAZ specimens.

For the weld metal Charpy specimens, the impact energies and ductile-to-brittle transition behaviour are more consistent than their corresponding HAZ results. For both the HYN and 101TM electrodes the impact energies show a consistently increasing trend of energy with temperature, as plotted in Figure 22. For the 101TM electrode there is a noticeable shift in the transition temperature associated with the high heat input welding conditions compared with the low heat input conditions. Table 9 summarizes the ductile-brittle transition temperatures for the weld metal Charpy tests.

The HAZ and weld metal impact energies associated with the GMAW panels are plotted together in Figure 23 with consistently increasing impact energy with increasing temperature.


0

25

50

75

100

125

150

175

200

225

250

-90 -80 -70 -60 -50 -40 -30 -20 -10 0 10 20Temperature (deg C)

Ener

gy (J

)HYN-LHHYN-HHBase Metal

a)

0

25

50

75

100

125

150

175

200

225

250

-90 -80 -70 -60 -50 -40 -30 -20 -10 0 10 20Temperature (deg C)

Ener

gy (J

)

101TM-LH101TM-HHBase Metal

b)

Figure 21 HAZ Charpy Impact Energy for the full sized specimens.


0

25

50

75

100

125

150

-90 -80 -70 -60 -50 -40 -30 -20 -10 0 10 20 30

Temperature (deg C)

Ener

gy (J

)

HYN-LHHYN-HH101TM-LH101TM-HH

Figure 22 Weld Charpy impact energies.

0

25

50

75

100

125

-90 -80 -70 -60 -50 -40 -30 -20 -10 0 10 20 30

Temperature (deg C)

Ener

gy (J

)

Weld MetalHAZ

Figure 23 Weld and HAZ Charpy Impact Energy from subsize (10x7.5 mm) samples.


Table 8 HAZ Ductile-Brittle Transition Temperatures assessed using the 50% Fracture Appearance Transition Temperature (FATT) and mid-shelf impact energy.

Base Plate HYN-LH HYN-HH 101TM-LH 101TM-HH GMAW 50% FATT

N/A -50 -5 -50 0 -30

Mid-Shelf -40 -50 0 -50 -5 -40

Table 9 Weld Metal Ductile-Brittle Transition Temperatures assessed using the 50% FATT and mid-shelf impact energy.

HYN-LH HYN-HH 101TM-LH 101TM-HH GMAW 50% FATT -45 -55 -60 -50 -30 Mid-Shelf -45 -50 -40 -40 -30 All of the weld metal Charpy impact specimens meet the requirements for an ABS grade 4Y400 [10], specified as 47J at -40C for the full size specimen and 31.3J at -40C for the subsize specimen.

3.6 Prior Austenite Grain Growth

Along with the cooling rate, another factor in determining the product of the austenite-ferrite phase transformation is the size of the prior-austenite grains [3]. In low carbon steels, coarser austenite grains may promote the formation of a harder transformation product (i.e. bainite) instead of ferrite-pearlite structure. Furthermore, the coarser the grain size, the coarser the resulting microstructure, resulting in wider spacing between bainite packets. Generally, the acicular ferrite is associated with good impact toughness, while lamellar structures including grain boundary ferrite and upper bainite are undesirable.

Austenite grain growth in the HAZ is a function of the complete thermal cycle, including the time above the grain-coarsening temperature. The grain coarsening temperature is influenced by the solubility kinetics of grain boundary pinning precipitates, such as TiN and AlN, and therefore is related to the micro alloying additions of the steel. Above the grain coarsening temperature, the grain boundary motion is un-impeded.

Grain sizes were measured from colour etched samples at 500 μm increments away from the straight fronted fusion line. The samples were etched using a Klemm-I reagent which colour tints ferrite different shades of blue and brown depending on the crystallographic orientations of the grains. This reagent was initially selected since it was believed to tint the prior-austenite grains rather than ferrite. However, since the ferrite grains transform from the prior austenite grains, fine ferrite grains may be associated with a finer austenitic structure.

Grain size measurements were made using an ASTM comparative approach [12] in which the colour etched grains were compared with a superimposed grid. At least 20 grain measurements were acquired along lines parallel to the fusion boundary. The average grain size and 95% confidence intervals for the four FCAW panels are plotted in Figure 24. Although the measured grain sizes correspond to the transformed product, rather than prior-austenite phase, grain sizes increase with decreasing distance away from the fusion line. In particular, the grain sizes


sampled by the low and high heat input heat affected zones are significantly different (See Figure 25).

0.0

20.0

40.0

60.0

80.0

100.0

120.0

0.00 0.50 1.00 1.50 2.00 2.50 3.00 3.50 4.00

Distance from Fusion Line (mm)

Ave

rage

Gra

in D

iam

eter

(μm

)

HYN LHHYN HH101TM LH101TM HH

Average Base

Approx. Charpy Notch Location

Figure 24 Average and 95% confidence intervals of prior-austenite grains diameters measured in

the heat affected zone. Data points offset for clarity.


a) HYN-HH 1.0 mm away from the fusion line. This region corresponds with the coarse grained zone (CGHAZ)

b) HYN-LH 1.0 mm away from fusion line. The region corresponds with the intercritical zone (ICHAZ)

Figure 25 Microstructures sampled by Charpy Impact Notches for the high heat input (a) and low heat inputs welds (b)


4 Discussion

HSLA-65 (ASTM A945) is a clean commercial micro-alloyed steel product with low carbon content for improved weldability, toughness and strength compared with conventional ABS grades DH/EH 36 [10]. The controlled rolled plate material used in this study resulted in a fine but banded ferrite and pearlite structure, which provides the requisite base plate toughness and strength.

During welding, a thermal gradient develops within the Heat Affected Zone (HAZ) that ranges from the fusion temperature to the sub-critical (i.e. less than Ac1). In the current study, the influence of the weld heat input varied the width of the various zones within the HAZ, the thermal gradient and cooling rates. For all regions in the HAZ, the welds produced using the higher heat input conditions resulted in longer dwell times at elevated temperatures and slower cooling rates compared with the lower heat input welds. The effect of the thermal cycle, dwell time and cooling rates on the subsequent microstructural development, hardness, strength and impact energy were investigated.

It is known that the HAZ properties are influenced by both the chemical composition and base metal microstructure [4]. In particular, the reduced HAZ toughness is partly related to grain growth and the resultant microstructure. It is generally recognized that the reduction in the carbon equivalent reduces the likelihood of high hardness untempered martensite formation, which is beneficial to improving toughness. For titanium treated micro-alloyed steel, such as HSLA-65, a fine dispersion of TiN precipitates has been found to be effective in pinning austenite grains up to 1400oC, thus retarding austenite grain coarsening, while decreasing the levels of impurities (Sulphur, Phosphorous, Nitrogen and Oxygen) within the base and HAZ increases the toughness. The simultaneous effect of the low carbon content and pinning effect of the TiN precipitates suppresses the formation of proeutectoid ferrite (grain boundary ferrite) and the grains transform completely to an acicular structure (upper bainite, lower bainite or martensite) [3, 4, 13]. The formation of the acicular phases is largely governed by the cooling rate, austenite grain size and presence of a fine dispersion of precipitates, including titanium oxides and nitrides and microalloying additions of Vanadium. To further investigate the influence of the HAZ thermal cycle on the resultant microstructures and in particular the cooling rate, a dilatometric study is currently underway for this steel.

The influence of the microstructure was most pronounced in the ductile-to-brittle transition behaviour since the bainitic microstructure forms a hard, high-strength phase as evidenced by the microhardness traverses and transverse tensile behaviour. Weldment hardness traverses across the weld showed limited softening behaviour and a smooth transition from the softer base to the harder weld metal. Similarly, all of the transverse tensile specimens failed away from the HAZ in the region of least constraint between the weld and the grip region.

The most pronounced influence of the welding process is the significant shift of the HAZ ductile-to-brittle transition behaviour to higher temperature. However, interpretation of the influence of heat input is confounded by the microstructure sampled by notch. For the HAZ notched specimens, this sampling variation accounts for the large scatter in impact energy and implied influence of the heat input.


The influence of the sampled microstructure was examined from a number of specimens in the ductile-to-brittle transition region (see Figure 26). The fracture surfaces of two specimens displaying upper-shelf (ductile) behaviour and three displaying lower-shelf (brittle) behaviour were sectioned along the specimen centerline in the direction of the crack propagation, thus revealing the relative location of the notch to the fusion boundary. Two representative specimens displaying either upper-shelf or lower-shelf behaviour are displayed in Figure 27 and Figure 28, respectively. For all the sectioned specimens, the notch root radius was located within 500 μm of the oscillating fusion boundary. The oscillating fusion boundary was caused by weaving the welding torch, which was necessary to retain arc and weld pool stability, especially at the low travel speeds used in these high heat input conditions. Considering this oscillation and the slight through-thickness variation of the fusion boundary, (see Figure 8a) the relatively consistent placement of the notch is impressive.

0

25

50

75

100

125

150

175

200

225

250

-90 -80 -70 -60 -50 -40 -30 -20 -10 0 10 20

Temperature (C)

Ener

gy (J

)

-5C160J

0 C60 J

Figure 26 Charpy Impact energy for HYN-HH notched 1.0 mm from the fusion line. Red markers indicated samples which were selected for sectioning while the labelled specimens are reported

herein and referred to as the upper and lower shelf specimens.

While the notch location is relatively consistent, the initial fracture paths and fracture morphology between specimens displaying upper and lower shelf behaviour is dramatically different. For the specimen displaying upper-shelf behaviour, the initial crack propagation involves significant ductile tearing as evident by the rough surface directly above the notch in Figure 27a and by the corresponding high number of voids adjacent to the notch in Figure 27b. Furthermore, the initial crack propagation was away from the fusion line into the intrinsically tougher base metal. In comparison, the fracture surfaces of the specimen displaying lower shelf behaviour had limited


regions of ductile tearing on either the fracture surface or cross section, with the initial crack propagation towards the fusion line and more brittle CGHAZ. In the ductile to brittle transition region, the fracture path of specimens displaying lower-shelf behaviour propagated along or close to the fusion boundary, while for those displaying upper-shelf bevaviour, the fracture path was significantly further away from the fusion line.

Therefore, from the evidence provided by the fracture paths, the variability in the ductile-to-brittle behaviour is directly related to the microstructural regions through which the crack propagates and, in particular, the crack initiation region associated with the notch. The scatter associated with the ductile-to-brittle transition region shown in Figure 26 is therefore intrinsically related to the fracture path within the microstructural gradient between the fusion boundary and the parent material. As a result of the intrinsic nature of the scatter, the ductile-to-brittle curve is drawn at the maximum temperature that provides the most conservative estimate of the transition behaviour.

Poor HAZ Charpy impact energy behaviour is attributed to the formation of an upper bainitic microstructure in which the aligned second phase facilitates the formation of cleavage facets. As shown in the fracture surface of one of the HYN-HH HAZ Charpy specimens in Figure 29, secondary cleavage cracks that run parallel to the principle fracture are visible in the relatively coarse bainitic structure, while the crack propagates along the aligned cementite constituent. The presence of these secondary cracks within the bainitic structure could be related to local brittle zones which are frequently cited in fracture toughness characterization of welded joints. Unlike the upper bainitic microstructure of the CGHAZ, the microstructure of all three weld metals was predominantly acicular ferrite. As a result of its fine ‘basket weave’ structure, there are no preferential paths for crack propagation and correspondingly the weld metal has a greater toughness than the CGHAZ. The suppression of the upper bainite transformation in the CGHAZ to acicular ferrite would undoubtedly improve the toughness.

The substantial difference in the impact transition behaviour for the low and high heat input HAZ of HSLA-65 is due in part to the microstructure sampled by the Charpy Notch and in part to intrinsic material differences. The microstructure sampled by the notch in the high heat input welds was located within the CGHAZ, while the notch in the low heat input welds was within the ICHAZ. Thus comparisons in impact energy made between these two heat inputs must reflect these microstructural differences.


a) Charpy Fracture Surface

b) Transverse cross section at line indicated in a) The crack propagation is from left to right.

Figure 27 Fracture surface of specimen displaying upper-shelf behaviour as highlighted in oFigure 26 (160J @-5 C and 60% cleavage appearance) b) transverse cross section along the

centerline. Arrows in a) and b) point to secondary cracks shown in Figure 29


a) Charpy Fracture Surface

b) Transverse cross section at line indicated in a)

Figure 28 Fracture surface of specimen displaying lower-shelf behaviour as highlighted in oFigure 26 (60J @0 C and 90% cleavage appearance) b) transverse cross section along the

centerline.


Figure 29 Micrograph of secondary cracks in the upper bainitic structure.

While the specified high heat input used in this trial is outside of those typically specified for flux core welding, it is easily obtained using higher heat input processes such as submerged arc welding (SAW). As was demonstrated in this welding trial, the heat input influences both the width of the HAZ, as well as the cooling rate and transformation products. The greater width of the HAZ in higher heat input welding processes results in a larger volume of material in which local brittle zones could be present. While the development of a microstructure susceptible to the creation of local brittle zones exists for both high and low heat input processes, the width of this region is significantly different. In addition, overmatched weld strength may provide a mechanism which shields the narrower local brittle zones adjacent to fusion line.

The characterization of the impact energy at a fixed distance away from the fusion line does not take into account the relative widths of the heat affected zones. Alternatively, the notches should be located in regions which experience similar thermal profiles and have a more consistent transformation product. However, such an approach would likely require a fatigue pre-cracked rather than blunt notch to ensure that the crack path sampled the intended region. In addition, the use of specimens with greater crack front constraint would better represent the structure. This can accomplished through the use of Drop Tower (DT) specimens, which have been found to further shift the brittle-ductile energy transition to higher temperatures.


5 Conclusions

The influences of welding heat input, welding process and filler metal composition were investigated for the joining of a micro-alloyed steel plate suitable for ship construction. During the welding of the plates, welding thermal histories were recorded adjacent to the fusion boundary and the Heat Affected Zone, along with the solidifying weld metal. The study examined the metallurgical and mechanical properties of the weldments, including tensile strength, hardness and impact energy. The major conclusions drawn from the investigation are that:

o For all cases, the filler metal and welding conditions produced welds which over-match the base metal strength. The weld metal had a fine acicular ferrite microstructure which exhibited good strength and a suitably low ductile-to-brittle transition temperature which exceed the requirements for an ABS grade 4Y400 consumable, making this a suitable pairing for this HSLA-65 steel.

o For the flux core welded panels, the most dramatic influence of the heat input was the significant increase in the ductile-to-brittle transition temperature with increasing heat input. The change in the heat input was associated with the relative width of the HAZ and microstructural region sampled by the relatively blunt notch.

Further characterization of the heat affected zone of this HSLA-65 steel is necessary in order to understand the substantial shift in the ductile-to-brittle transition behaviour with increasing weld heat input.




Annex A FCAW Plate Usage

A.1 71T1HYN-Low Heat Input Panels

30 As received panel showing the relative location of the three blocks labelled A, B and C.Figure


855 mm

8X WM DT 9X FL DTTTAT1

AW1 AF1AW2 AW3 AW4 AW5 AW6 AW7 AW8 AF2 AF3 AF4 AF5 AF6 AF7 AF8 AF9

Figure 31 Specimen locations and labels from Block A of panel 71T1HYN-LH.

AW1

AW2

AF1AF2

310.00 mm

120.00 mm

Figure 32 Specimen locations and labels from block B of panel 71T1HYN-LH.


4X AWM TensileLabel AT1 to AT4

710

17X WM Charpy 17X FL Charpy

Label AW1 to AW17 Label AF1 to AF17

Figure 33 Sample locations and labels from Block C of panel 71T1HYN-LH.

A.2 71T1HYN-High Heat Input Panels

Weld Plate HYN HHTop and Bottom Views After CuttingMay 10, 2007

2nd WTHYN HH-A

Weld

1st WT 1st WT

HYN HH-B 2nd WT HYN HH-C

HYN HH-CEND

HYN HH-B

Thermocouples

3 2 X 1 1011 9

ENDHYN HH-A

IP REPAIRWeld

817625

148

310 1313

396

14 135232

332

61333

613396

315490

817

146

310 10

HYN HH-B3X Expansion of Thermocouple Positions

2.8 mm7.4 mm3.2 mm1.5 mm

35 mm

87 mm111 mm

61 mm

6.4 mm3.2 mm

2.2 mm

328 mm354 mm

371 mm

Weld3 2 X 1 10 11 9

Figure 34 As received panel showing the relative location of the three blocks labelled A, B and C.


817 mm

8X WM DT 8X FL DTTTBT1

BW1 BF1BW2 BW3 BW4 BW5 BW6 BW7 BW8 BF2 BF3 BF4 BF5 BF6 BF7 BF8

Figure 35 Specimen locations and labels from Block A for panel 71T1HYN-HH

310.00 mm

BW1 BW2 BW3 BW4BF1 BF2 BF3

215.00 mm

Figure 36 Specimen locations and labels from Block B for panel 71T1HYN-HH


3X AWM TensileLabel BT1 to BT3

613


Label BW1 to BW16 Label BF1 to BF16

Figure 37 Specimen locations and labels from Block C for panel 71T1HYN-HH

A.3 101TM -Low Heat Input Panels Bottom View

Weld Plate 101TM LHTop and Bottom Views After CuttingMay 10, 2007

101TM LH-B

Thermocouples

1110 9 3 2 1

273

Weld

END101TM LH-A

10

137

275

850

101TM LH-CEND

705Note: Weld width varied from10 mm near ends to 6 mmnear middle

850743

101TM LH-A129

820

6275

Weld

101TM LH-B

273

62216Top View

101TM LH-C

705

101TM LH-B3X Expansion of Thermocouple Positions

11 10 9 3 2 1

Weld10.0 mm 5.7 mm 3.9 mm 10.2 mm 6.6 mm 4.4 mm

25 mm

51 mm77 mm

197 mm222 mm

248 mm

6 mm

Note: Vertical distances are from theedge of 6 mm wide weld



850 mm

8X WM DT 9X FL DTTTCT1

CW1 CF1CW2 CW3 CW4 CW5 CW6 CW7 CW8 CF2 CF3 CF4 CF5 CF6 CF7 CF8 CF9

Figure 39 Specimen locations and labels from Block A for panel 101TM-LH

CW1

CW2

CF1CF2

310.00 mm

120.00 mm

Figure 40 Specimen locations and labels from Block B for panel 101TM-LH


4X AWM TensileLabel CT1 to CT4

705


Label CW1 to CW17 Label CF1 to CF17

Figure 41 Specimen locations and labels from Block C for panel 101TM-LH

A.4 101TM –High Heat Input Panels

Weld Plate 101TM HHTop and Bottom Views After CuttingMay 10, 2007

Weld

END101TM HH-A

855

144

15310

Thermocouples

101TM HH-B

11 10 9 3 2 1

271

END101TM HH-C

702

101TM HH-C

SLAG SPILL

702

397302

101TM HH-B

WT WT

271

37217

Weld

101TM HH-A

855

150

1113

310

101TM HH-B3X Expansion of Thermocouple Positions

11 10 9 3 2 1Weld7.9 mm 3.4 mm 2.5 mm 8.4 mm 4.4 mm 2.5 mm

24 mm

50 mm76 mm

191 mm218 mm

242 mm

Bottom View

Top View



855 mm

8X WM DT 9X FL DTTTDT1

DW1 DF1DW2 DW3 DW4 DW5 DW6 DW7 DW8 DF2 DF3 DF4 DF5 DF6 DF7 DF8 DF9

Figure 43 Specimen locations and labels from Block A for panel 101TM-HH

DW1

DW2

DF1DF2

310.00 mm

120.00 mm

Figure 44 Specimen locations and labels from Block B for panel 101TM-HH


4X AWM TensileLabel DT1 to DT4

702


Label DW1 to DW17 Label DF1 to DF17

Figure 45 Specimen locations and labels from Block C for panel 101TM-HH




Annex B GMAW Plate Usage

Discard

Transverse Tensile

T BCharpyCharpy

Discard first 20 mm of weld Discard first 15 mm of Plate

224 mm 15 mm

313 mm

228 mm

AWM Tensile 1

AWM Tensile 2

AWM Tensile 3

20 mm

Figure 46 GMAW Plate 1. The central region contained the drilled holes for the thermocouples.


Tran

sver

se T

ensil

e

T B

17 FusionLine Charpy

4 FusionLine DT

350

20 222 20830

Disc

ard

Disc

ard

Figure 47 GMAW Plate 2 showing the location of various specimens.


3 FusionLine DT

T B

350

Disc

ard

Disc

ard

2 WM DT 5 WM DT

25 215 21325

Figure 48 GMAW Panel 3 showing the location of various specimens.




Annex C Weld Thermal Monitoring

C.1 GMAW Plate 2

0

100

200

300

400

500

600

700

0 10 20 30 40 50 6

Time (s)

Tem

pera

ture

(C)

T3

T2

T1

0

Figure 49 First Pass – GMAW Plate 1


0

100

200

300

400

500

600

700

800

0 10 20 30 40 50 6

Time (s)

Tem

pera

ture

(C)

T3

T2

T1

0

Figure 50 Second Pass- GMAW Plate 2


0

100

200

300

400

500

600

700

800

900

1000

0 10 20 30 40 50 6Time (s)

Tem

pera

ture

(C)

0

T1

T2

T3

Figure 51 Third Pass – GMAW Plate 2


C.2 GMAW – Plate 3

0

100

200

300

400

500

600

0 10 20 30 40 50

Time (s)

Tem

pera

ture

(C)

60

T1

T2

T3

Figure 52 First Pass – GMAW Plate 3


0

100

200

300

400

500

600

700

0 10 20 30 40 50

Time (s)

Tem

pera

ture

(C)

60

T1

T2

T3

Figure 53 Second Pass – GMAW Plate 3


0

100

200

300

400

500

600

700

800

900

0 10 20 30 40 50

Time (s)

Tem

pera

ture

(C)

60

T1 (T8-5=7.6 s)

T2

T3

Figure 54 Third Pass – GMAW Plate 3


0

100

200

300

400

500

600

700

800

900

1000

0 10 20 30 40 50

Time (s)

Tem

pera

ture

(C)

60

T1

Figure 55 Fourth Pass – GMAW Plate 3


References .

1. ASTM, A945/A945M Standard Specification for High-Strength Low-Alloy Structural Steel Plate with Low Carbon and Restricted Sulfur for Improved Weldability, Formability, and Toughness, in Iron and Steel Products. 2005, ASTM.

2. Posada, M.D., J. Characterization of HAZ properties and heat-treatment and TMCP HSLA-65 Steels. in Thermec 2000. 2000. Las Vagas, NV.

3. Lancaster, J.F., Metallurgy of Welding. 6 ed. 1999: Abington. 446.

4. Stern, I.L.W., M.; Ku, D.Y, Higher strength steels specially processed for high heat input welding. Journal of Ship Production, 1985. 1(4): p. 222-237.

5. Sampath, K., An Understanding of HSLA-65 Plate Steels. Journal of Materials Engineering and Performance, 2006. 15: p. 32.

6. Bailey, N., Weldability of ferritic steels. 1994: Woodhead Publishing.

7. Begg, D.,Re: Weld Inclusions in HSLA-65 Welds Email,2007

8. Geels, K., Metallographic and materialographic specimen preperation, light microscopy, image analysis and hardness testing. 2007, Lancaster, PA: ASTM.

9. Dexter, R.J. and M. Ferrell, Optimum weld-metal strength for high-strength steel structures. 1995, Ship Structures Committee: Bethlehem. p. 132.

10. ABS, Rules for Materials and Welding. 2008. p. 383.

11. ASTM, E23-02, Standard Test Methods for Notched Bar Impact Testing of Metallic Materials, 3.01, American Society for Testing Materials, 2003

12. ASTM, E112 Standard test method for determining average grain size, in Metals test methods and Analytical Procedures. 1996, ASTM.

13. Gladman, T., The Physical Metallurgy of Microalloyed steels. 1997: The institute of materials.



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1. ORIGINATOR (The name and address of the organization preparing the document. Organizations for whom the document was prepared, e.g. Centre sponsoring a contractor's report, or tasking agency, are entered in section 8.) Defence R&D Canada – Atlantic 9 Grove Street P.O. Box 1012 Dartmouth, Nova Scotia B2Y 3Z7

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3. TITLE (The complete document title as indicated on the title page. Its classification should be indicated by the appropriate abbreviation (S, C or U) in parentheses after the title.) The Influence of Heat Input on the Fracture and Metallurgical Properties of HSLA-65 Steel Welds: Fabrication conditions, Tensile, Impact and Microstructures

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HSLA-65 steel (ASTM A 945-05) is regarded as an excellent naval ship steel. Its low carbon equivalent and strength-to-weight ratio make it a candidate for future Canadian naval platforms. Unlike previous naval platforms, future ships might be expected to be deployed to Arctic waters. This increased operational requirement imposes a strict fracture toughness criterion for both the base metal and welds. In particular, the heat affected zone near the weld is of concern, especially for high heat input processes. In this study, welds were fabricated using the flux core arc welding and Gas Metal Arc processes with variations in the welding heat input and welding consumable. While all of the welded panels exhibited satisfactory tensile and hardness properties, a dramatic shift in the ductile-to-brittle transition temperature to higher temperatures was observed for high heat input conditions when the notch was positioned within the heat affected zone. The interpretation of this shift in transition behaviour was confounded by the relative width of the coarse grained heat affected zone, the microstructure sampled by the advancing crack and the position of the blunt Charpy notch.

14. KEYWORDS, DESCRIPTORS or IDENTIFIERS (Technically meaningful terms or short phrases that characterize a document and could be helpful in cataloguing the document. They should be selected so that no security classification is required. Identifiers, such as equipment model designation, trade name, military project code name, geographic location may also be included. If possible keywords should be selected from a published thesaurus, e.g. Thesaurus of Engineering and Scientific Terms (TEST) and that thesaurus identified. If it is not possible to select indexing terms which are Unclassified, the classification of each should be indicated as with the title.) HSLA-65; Welds; Mechanical Properties; HAZ


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