2 7
AüS. 1968
FRWWISCH5 HoGFSCHOOL OELR
VUEGTUIGBOUWKUNDE
BIBLIOTHEEK
CoA R E P O R T MAT No. 2
THE COLLEGE OF AERONAUTICS
CRANFIELD
THE WELD HEAT A F F E C T E D ZONE STRUCTURE
AND P R O P E R T I E S O F TWO MILD S T E E L S
by
CoA REPORT MAT. No. 2 January, 1968
• THE COKLEGE OF AERONAUTICS CRAETTELD
The veld heat affected zone structure and properties of two mild steels
by
-E. Smith, Ph.D., B . S c , A.I.M., M.D. Coward, M.Phil., B . S c , A.R.S.M., and R.L. Apps, Ph.D., B . S c , F.I.M., A.M.Inst.W.
CORRIGEl^A
Fig. 5 to read: Temperature of 788°C. Fig. 8 to read: Magnification x 7OO Fig. 12 to read: x 150OO not x 15OO Fig. 12 to read: x l4000 not x lUOO
V
CoA R E P O R T MAT No. 2 J a n u a r y , 1968.
THE COLLEGE O F AERONAUTICS CRANFIELD
T h e Weid Heat Affected Zone S t r u c t u r e and P r o p e r t i e s of Two Mild S t e e l s
by
E . S m i t h , P h . D . . B . S c , A . I . M . , M . D . C o w a r d , M . P h i l . . B . S c , A . R . S . M . ,
SUMMARY
T h e v i s i b l e weld heat affected zone (HAZ) of mild s t e e l i s shown to be m a d e up of four d i s t i n c t r e g i o n s (a) the r e g i o n of g r a i n c o a r s e n i n g , (b) the r e g i o n of g r a i n r e f i n e m e n t , (c) the p a r t i a l l y t r a n s f o r m e d r e g i o n , and (d) the r e g i o n of s p h e r o i d i s a t i o n . R e p r o d u c t i o n of t h e s e HAZ s t r u c t u r e s in un-notched C h a r p y s p e c i m e n s , 1.0 c m . s q u a r e and 7 . 5 c m . long, h a s been achieved u s i n g a c o n t r o l l e d r e s i s t a n c e heating a p p a r a t u s capable of cycling s p e c i m e n s through t h e r m a l c y c l e s m e a s u r e d in the a c t u a l HAZ of a s u b m e r g e d a r c b e a d - o n - p l a t e weld in I j " thick plate using a h e a t input of 108 k i l o j o u l e s / i n c h .
M e t a l l u r g i c a l examination and h a r d n e s s m e a s u r e m e n t s on two different mild s t e e l s have shown that the s i m u l a t e d r e g i o n s a r e c o m p a r a b l e with the a c t u a l r e g i o n s in the weld H A Z . The m e c h a n i c a l p r o p e r t i e s of the s i m u l a t e d HAZ r e g i o n s have been m e a s u r e d and c o r r e l a t e d with the m e t a l l u r g i c a l m i c r o -s t r u c t u r e -s a -s d e t e r m i n e d by m e a n -s of optical and e l e c t r o n m i c r o -s c o p y .
The r e s u l t s show that a m a r k e d e m b r i t t l e m e n t o c c u r s in both s t e e l s in the r e g i o n of g r a i n c o a r s e n i n g and t h i s i s a s s o c i a t e d with a v e r y c o a r s e
Widmanstatten s t r u c t u r e . T h i s i s p a r t i c u l a r l y d e t r i m e n t a l t o the s t r e n g t h of the welded joint s i n c e t h i s r e g i o n contains the toe of the weld which m a y act a s a c o n s i d e r a b l e s t r e s s - r a i s e r , and r e s i d u a l s t r e s s e s a r e often high in t h i s r e g i o n .
A s i l i c o n killed s t e e l t o BS1501-161 G r a d e B i s shown to be e m b r i t t l e d in all r e g i o n s e x p e r i e n c i n g a peak t e m p e r a t u r e above to A c , . Between t h e A c , and t h e Ac„ p o i n t s , t h i s is a t t r i b u t e d in p a r t to the f o r m a t i o n of u p p e r b a i n i t e . A p a r t i a l l y deoxidised s t e e l to BS15 is shown t o have i m p r o v e d notch t o u g h n e s s p r o p e r t i e s in the r e g i o n of t h e HAZ o u t s i d e the r e g i o n of g r a i n c o a r s e n i n g and t h i s i s a s s o c i a t e d with a r e d u c t i o n of the f e r r i t e g r a i n s i z e .
A r e s t r a i n t applied d u r i n g t h e t h e r m a l c y c l e s which i m p o s e d a p e r m -anent d e f o r m a t i o n of 2 - 4% is shown to have no significant effect on t h e r e s u l t i n g m i c r o e t r u c t u r e and p r o p e r t i e s in BS15.
F i n a l l y a p o s t - w e l d heat t r e a t m e n t at about 650 C is shown to have no significant effect on t h e impact p r o p e r t i e s of BS1501.
CONTENTS Page 1. Introduction 1 2. Experimental 2 2.1 Materials 2 2.2 Experimental Procedure 3
2 . 2 . 1 Preparation and examination of welds 3
2 . 2 . 2 Hardness of the weld HAZ 4 2 . 2 . 3 Simulation of weld HAZ structures 4
2 . 2 . 4 Mechanical properties of the simulated weld HAZ
structures 4 2 . 2 . 5 The effect of post-weld heat treatment 5
3. Results 5 3.1 Metallography and Hardness of the Weld HAZ 5
3.2 Metallography of the Simulated Weld HAZ Structures 7
3.2.1 BS15 7 3.2.2 BS1501 7 3.3 Mechanical Properties of the Similated HAZ 8
3.4 The Effect of Post-Weld Heat Treatment 8
4. Discussion 8 4.1 Simulation of the Thermal Cycles 11
4.2 Hardness of the Actual Weld HAZ 11 4.3 Banding of the Parent Plate Structures 11 4.4 Mechanical Properties and Microstructures in BS15
Simulated Weld HAZ Regions 12 4 . 5 Mechanical Properties and Microstructures in BS1501
Simulated Weld HAZ Regions 14 4.6 The Effect of Restraint on the Simulated HAZ Properties
of BS15 15 4.7 The Effect of Post-Weld Heat Treatment 16
5. Conclusions 16 6. Bibliography 18
1
-1. Introduction
Mild steel is widely used in welded fabrications. Tests carried out on welded joints show that, in general, weldments in mild steel a r e acceptable, but that certain regions show a deterioration in mechanical properties which have a significant effect on the overall joint performances in certain c i r c u m -stances.
F r o m the weld fusion line to the outside edge of the visible weld HAZ the peak t e m p e r a t u r e s reached during the welding cycle vary by several hundred d e g r e e s . As a result this region, though usually only a few m i l l i m e t r e s wide, exhibits a wide range of naicrostructures and these a r e r e s p o n -sible for the variations in mechanical properties encountered in this region. These variations a r e produced by the temperature and strain cycles associated with the welding p r o c e s s .
There i s , therefore, an important need for the metallurgical s t r u c t u r e s and mechanical properties of the weld HAZ to be fully investigated. The practical difficulties associated with testing actual weld HAZ s t r u c t u r e s , due to the wide variety of structures occurring over a very narrow region and accurate positioning of test specimens has led to the development of a
simulation technique whereby the HAZ s t r u c t u r e s a r e reproduced in specimens large enough for mechanical testing by subjecting them to actual weld HAZ t h e r m a l cycles.
The aim of the present work was to examine the relationship between the mechanical properties and the m i c r o s t r u c t u r e s produced in the visible weld HAZ of two commercial mild steels for a typical set of welding conditions corresponding to a heat input of 108 kilojoules per inch. Bead-on-plate welds were made on both steels using an automatic submerged ar c welding machine and the range of m i c r o s t r u c t u r e s produced in the HAZ were examined using optical and electron miicroscopic techniques. The hardness values of the various m i c r o s t r u c t u r e s were also measured.
The r e s u l t s show that for both steels the visible weld HAZ can be divided into a number of distinct regions. These a r e the region of grain coarsening, the region of grain refinemient, the region of partial t r a n s f o r m -ation, and the region of spheroidization. This is in general agreement with the r e s u l t s of other w o r k e r s ^ ' ^ . although the region of spheroidization is such a narrow one, that not all workers have classified it as a distinct region3"5. However a region of spheroidization, approximately 0.1 mm. wide, was observed in the present s t e e l s .
The mechanical properties of the various regions of the weld HAZ have been determined on specimens of a suitable size for mechanical testing after treatment in an apparatus capable of reproducing the thermal cycles which were measured at various points in the HAZ during the preparation of a bead-on-plate weld.
The technique used for measuring the weld thermal cycles and results obtained have been reported previously*'. Four thermal cycles were chosen for the present work, and specimens of both s t e e l s , 1.0 cm. square and
2
-7.5 cm. long, have been subjected to these thermal cycles using the simulation equipment developed at Cranfield and described in detail elsewhere^. These 'simulated' specimens have been examined by optical and electron microscopy and the structures compared with those observed in the HAZ of the actual welds. Hardness, notch-impact, and tensile tests have been carried out on the simulated specimens to determine the relationships between the mechanical properties and the microstructures.
In addition, the effect of a heavy restraint on the specimen during the simulation of the thermal cycles has been investigated.
In general, the material affected by the weld thermal cycles had inferior notch toughness properties, lower ductility, and increased strength and hardness as compared to the parent plate material. The poorest
properties were associated with the region of grain coarsening. Restraint applied during the simulation work had no measurable effect on the resulting mechanical properties and m i c r o s t r u c t u r e s .
3 8
Some workers ' have reported the formation of martensite in former pearlitic areas where transformation had occurred from an austenite-ferrite microstructure during the welding of mild steel. The present work indicates that, under welding conditions of high heat input, this is not likely to occur.
Finally, the effect of a post-weld heat treatment on the weld HAZ properties of one of the steels has been investigated. The results indicate that such a treatment had only a very slight beneficial effect on the fracture toughness properties of the visible weld HAZ structures and, support the generally accepted principle that mild steel can normally be welded without a post-weld heat treatment except, possibly, in the case of very thick sections.
2. Experimental 2.1 Materials
One steel was a silicon-killed mild steel to BS1501-161 Grade B and was supplied as 2" thick hot rolled plate. The other was a mild steel to BS15 and was supplied in Ij" thick hot rolled plate. This steel was only slightly deoxidised and could be classified as intermediate between a rimming and a semi-killed steel. The chemical analysis of both materials is shown in Table 1, together with the mechanical properties of the BS1501 parent plate. No mechanical property data were supplied with the BS15 steel,
3
-TABLE I Chemical Analysis and Mechanical P r o p e r t i e s of BS1501 and BS15 BS1501 BSl.'i C .19 .21 Si . 2 3 .065 Mn . 6 4 .89 S .046 .050 P .032 .040 Ni .14 * n . d . C r . 1 2 n . d . a y 18.1 t . s . i . -a max 29.2 t . s . i . -E l . 29% -R.ofA. 55% -* n . d . - not determined. 2.2 Experimental Procedure
2 . 2 , 1 Preparation and examination of welds
One piece, 7" x 12". was flame cut from the BS1501 plate, and one piece, 26" X 12". was flame cut from the BS15 plate. The faces of each piece were then cleaned using a portable grinder and a bead-on-plate weld wfis made down the centre of each using a British Oxygen Company automatic
3 "
submerged a rc welding unit with r ^ mild steel filler wire and a heat input of 108 kilojoules per inch. The welding conditions identical with those used for tiie determination of the thermal cycles^ a r e summarised in Table II.
TABLE II Bead-on-plate welding conditions
Heat Input Arc Voltage Arc Current Welding Speed 108±2 Kilojoules/inch 30±2 Volts 390±10 Amps 6 | ± 5 Inches/min
After cooling to room temperature, the welds were sectioned and the sections prepared for m i c r o s c o p i c examination using normal grinding and polishing techniques. They were then etched in 2% nital for a period varying from 5 to 20 seconds, depending on the structure involved. Each section was then examined carefully using a Leitz microscope and a Reichert microscope to determine the range and extent of the microstructures produced in the weld HAZ.
The form and distribution of the carbides were then studied in more detail, up to magnifications of x 40,000. using the carbon replication technique and a Siemens electron microscope (model Elmiskop la) with a filament
4
-voltage of 2.5 volts and a beam current of 5-40 microamps. 2 . 2 . 2 Hardness of the weld HAZ
The variations in hardness across the HAZ of both welds were d e t e r m -ined by making hardness measurements at | m m . intervals, starting at the fusion boundary and finishing in the unaffected parent plate, A Zwick hard-n e s s testihard-ng machihard-ne, with a load of 5 kgms was used ahard-nd a mihard-nimum of 5 hardness measurements were made at each interval in order to reduce scatter in the re s ults ,
2 . 2 . 3 Simulation of weld HAZ structures
Because of the inherent difficulties in measuing the mechanical proper-ties of the various regions of the weld HAZ. a technique has been developed whereby specimens are subjected to thermal cycles similar to those experienced in the HAZ of a weld . The equipment built at Cranfield is designed to take 1cm, square specimens through any desired thermal cycle. Thermal cycles have been measured on the BS15 plate using the set-up described earlier in this report for a bead-on-plate weld with a heat input of 108 kilojoules per inch. The details of the temperature measurement and the development of the thermal simulation equipment have been described in detail"."^. Heating is carried out by self-resistance, and cooling by means of water cooled copper clamps. This equipment has been found to be capable of accurately simulating weld thermal cycles,
From the temperature measurement work four thermal cycles were chosen for simulation in each material, as representing particular regions of the weld HAZ, F o r the BS1501 material thermal cycles with peak temperatures of 788°C, 8930C. 1088°C and 1347°C were chosen, and for the BS15 material, thermal cycles with peak temperatures of 788°C. 893°C. 1070°C and 1305°C were chosen,
In the main part of the work, the specimens were held in a pair of moveable jaws, so that they could expand and contract freely during the thermal cycles. Another s e r i e s of t e s t s were carried out on the BS15 material in which the specimens were held rigidly in the jaws so that they could not expand or contract during the thermal cycles. The object of this was to determine the effect of restraint on the resulting microstructures and mechanical properties,
After the simulation treatment was completed, one specimen from each category was sectioned and examined optically; in addition carbon replicas were prepared and examined using the electron microscope. The structures were then compared with those formed in the HAZ of the actual weld,
2 . 2 . 4 Mechanical properties of the simulated weld HAZ structures Hardness tests were made on the simulated specimens, using the same testing conditions as before, and the results compared with the hardness values measured in the actual welds.
5
-The mechanical properties were further examined by machining the simulated blanks into standard Charpy V-notch specimiens and modified Hounsfleld number 13 tensile test pieces. The modification to the standard Hounsfleld tensile test piece became necessary as the volume of heat treated zone In the simulated specimens was too small to contain the gauge length of the standard test piece. Some preliminary tests were carried out on test pieces with a shortened gauge length, but In other respects Identical In dimensions to the Hounsfleld number 13 test piece, and it was found that a specimen with a gauge length of 0 . 3 " gave Identical tensile properties to the standard test plece^",
The tensile t e s t s were performed on a 10,000 lb, Instron Universal
testing machine at room temperature using a cross head speed of ,05 Inch/mln. The yield s t r e s s (or 0,2% proof s t r e s s ) , ultimate s t r e s s , elongation and
reduction of area were recorded,
The Charpy tests were carried out on a Losenhausen Impact testing machine with a striking energy of 217 f t , l b s . The t e s t s were carried out over a wide range of temperature and the fracture energy and percentage crystallinlty of the fractured faces recorded. The specimens were soaked In an appropriate constant temperature bath for 10 minutes prior to testing, Temperatures below ambient were obtained using a mixture of methyl alcohol and solid carbon dioxide; temperatures above ambient were obtained with a heated oil bath. F r o m these results the transition temperatures of the various structures were determined.
2 , 2 . 5 The effect of post-weld heat treatment
The effect of post-weld heat treatment was investigated by heating specimens of the BS15 steel, which had previously been thermally cycled In the simulator to peak temperatures of 788°C, 893°C and 1070OC, for 30 minutes at 650±50°C In an electric hump furnace. Temperature control of the furnace was very e r r a t i c during this heat treatment and consequently the temperature can only be given accurately to ±50°C. Charpy V-notch impact t e s t s were then carried out and the results compared with those obtained without the post-weld heat treatment.
3, Results
3,1 Metallography and Hardness of the Weld HAZ
The parent plate microstructures were typical of hot rolled structures and consisted of banded pearlite in a ferrite matrix, as shown In Figs, 1 and 9 for BS15 and In F i g s . 2 and 14 for BS1501. BS15 had a much c o a r s e r structure than BS1501. The changes in the parent plate microstructures due to the bead-on-plate welds are shown in F i g s . 1 and 2. The HAZ could be divided Into four distinct regions and these were Identified in t e r m s of their distance from the weld fusion boundary In BS15 as
follows; 6 follows;
-(a) The region of grain coarsening
This region extended from 0 to 1.2 mm. from the fusion boundary and experienced maximum tentiperatures In the approximate range 1500°C to
1100°C. Since the maximum temperatures were well Into the austenltlc region, austenite grain growth took place and the grain size diminished with Increasing distance from the fusion boundary and corresponding decreasing maximum t e m p e r a t u r e s . In the neighbourhood of the fusion zone the relatively rapid cooling rates were evidenced by the Widmanstatten ferrite structure, as shown In F i g s . 1(a) and 2(a),
(b) The region of grain refinement
This region extended from 1.2 to 3.0 mm. from the fusion boundary and experienced maximum temperatures in the approximate range 1100°C to 900°C, This region was heated to just above the temperatures required for complete austenltization and exhibited a totally refined equlaxed grain structure, as shown in F i g s , 1(b) and 2(b),
(c) The region of partial transformation
This region extended from 3,0mm. to 4,5mm, from the fusion boundary and experienced maximum temperatures in the approximate range 900°C to 750°C. This region was heated between the temperature limits of Ac. and Ac which produced partial austenltization of the original structure. The regions richest in carbon, where parial dissolution of the pearlite had occurred, exhibited extremely fine grained s t r u c t u r e s , whilst the ferrite regions became less affected as the maximum temperatures attained diminished and tended to approach the original grain size. No t r a c e of martensite formation was observed In this region although a transformation product resembling balnlte was observed in the BS1501. These structures a r e shown in F i g s , 1(c) and 2(c).
(d) The region of spheroidisation
Over a narrow region, approximately 4 , 5 m m . to 4.6mm. from the fusion boundary, heated to maximum temperatures just below the Ac.
(~ 750°C), the lamellar carbides tended to dissolve and reform as spheroidal particles upon cooling. The ferrite matrix remained unaffected by the
relatively low maximum temperatures reached in this region and the very short time when the temperature was near the A c . , Fig, 1(d) shows a pearlite grain In the BS15 steel which had been spheroldlsed (marked by letter B) adjacent to one that had just partially transformed (marked by letter A), Fig. 2(d) shows the corresponding region In the BS1501 steel,
Beyond this region structural changes were not observed by optical microscopy and, therefore, this was defined as the boundary of the visible weld HAZ,
The changes in hardness through the weld HAZ with increasing distance from the fusion boundary are shown In F i g s , 3 and 4. In general, it can be
hardness, and this decreased with Increasing distance from the fusion boundary. This effect was more marked In BS1501,
3.2 Metallography of the Simulated Weld HAZ Structures 3.2.1 B.S.15
(a) Simulated to a peak temperature of 788°C
The structure produced by this treatment Is shown In Fig, 5; it is similar to the partially transformed region shown In Fig, 1(c). but there is some evidence of spheroidisation. Feathery upper bainite is observed in the electron micrographs so it is likely that the peak temperatures experienced have been just above the Ac. ,
(b) Simulated to a peak temperature of 893°C
The structure produced by this treatment is shown in Fig, 6. Grain refinement of the ferrite Is apparent and aclcular upper bainite has formed In the carbon-rich a r e a s Indicating that peak temperatures have been above Ac„, This appears to be equivalent to the region of grain refinement,
Ó
(c) Simulated to a peak temperature of 1070°C
The structure produced by this treatment Is shown in Fig, 7, A Widmanstatten distribution of ferrite can be observed and the grain size Is much c o a r s e r than In the specimen cycled to 893°C, This suggests that this is the start of the region of grain coarsening,
(d) Simulated to a peak temperature of 1305°C
The structure produced by this treatment is shown In Fig. 8, A very coarse network of Widmanstatten ferrite has been produced, suggesting that this region r e p r e s e n t s the area immediately adjacent to the fusion boundary,
3 . 2 . 2 B . S . l S O l
(a) Simulated to a peak temperature of 788°C
The structure produced by this treatment is shown in Fig, 10, There has been a refinement of the ferrite grain size. Large pools of carbide resembling feathery upper balnlte have formed In the former pearlitic a r e a s . suggesting that peak temperatures have been above the Ac. . This structure Is therefore considered to represent the region of partial transformation,
(b) Simulated to a peak temperature of 893°C
The structure produced by this treatment is shown in Fig, 11, This is similar to the previous sample, although the ferrite grain size is now s o m e -what larger and there appears to have been more homogenlsation of the carbon. The structure suggests that peak temperatures have been approaching the Ac .
(c) Simulated to a peak temperature of 1088 C
The structure produced by this treatment is shown in F i g . 12, A coarse Widmanstatten structure of proeutectoid ferrite and upper balnlte has been produced, suggesting that the peak temperature has been well in excess of the Ac„ and into the region of grain coarsening
(d) Simulated to a peak temperature of 1347 C
The structure produced by this treatment is shown in Fig, 13, A very coarse grained Widmanstatten structure of proeutectoid ferrite and upper bainite has been produced. Indicative of peak temperatures close to the melting point,
3,3 Mechanical Properties of the Simulated HAZ
The hardness of both materials simulated through the four thermal cycles a r e recorded with the micrographs. F i g s , 5-14,
The results of the Charpy V-notch impact tests on specimens of both m a t e r i a l s simulated through four thermal cycles representing different parts of the weld HAZ a r e shown in F i g s , 15-18,
The amount of plastic deformation produced in the specimens during thermal cycling under the condition of full restraint Is obtained from the overall displacement of the copper blocks during cycling. Fig, 19 shows displacement-time curves under zero and full restraint for a B,S,15 specimen where cycling reached a peak temperature of 1070°C, The difference In displacement between the curves for zero restraint and for full restraint after approximately 130 seconds represents the amount of permanent deformation produced by the restraining mechanism,
The percentage plastic deformation Introduced In the full-restraint
condition for the three thermal cycles Is 2% for a peak temperature of 788°C, 3% for 8930C. and 4% for 1070°C,
The results of the tensile tests on the BS1501 steel, using the Hounsfleld number 13 test pieces with a modified gauge length of 0 . 3 " . are illustrated In Fig, 20 and summarised In Table III, The values of transition temperature,
calculated from the impact r e s u l t s , and the hardness values are also included in Table III, (See page 9),
3,4 The Effect of Post-Weld Heat Treatment
The effect of a post-weld heat treatment, in which specimens of the BS15 steel were heated for 30 mins, at 650°C±50 C after simulation to 788°C, 893°C and 1070°C is summarised In F i g s , 21 and 22. Fig, 21 shows the effect of the treatment of the energy required for fracture In the Charpy V-notch Impact t e s t , and Fig, 22 shows the corresponding fracture appearance transitions,
4 . Discussion
One of the main difficulties In attempting to explain the structural changes occurring during the welding process is that, due to the marked departure from equilibrium conditions, caused by the rapid heating and cooling rates and the short times at peak t e m p e r a t u r e s , conventional equilibrium data are not
TABLE lU Mechanical Properties of BS1501 mild steel simulated specimens
Treatment As received Cycled to 788°C Cycled to 893°C Cycled to 1088°C Cycled to 1347°C 0.2% Proof s t r e s s ( t . s , i . ) 18.9 20.3 22.6 26.4 33.5 U , T , S , ( t , s . l , ) 30.4 33.9 38.3 39,0 43.3 Reduction of Area (%) 79 72 67 74 68 Elongation (%) 59.3 54.3 53.0 52.8 48,0 U . T . S . Proof s t r e s s 1.61 1.67 1,70 1,49 1.29 Hardness (HVIO) 138 170 190 207 231 Transition temp, at 50% Crysty -7°C 39°C 21°C 27°C 43°C Transition temp, at 20 f t , l b s . -23°C 4 ° C -16°C
o°c
5°C 1
COTable IV summiarises the values of transition temperature and hardness for the BS15 steel.
TABLE IV Mechanical Properties of BS15 mild steel simulated specimens
Treatment As received Cycled to 788°C Cycled to 893°C Cycled to 1070°C Cycled to 1305°C Hardness HV 5 199 200 205 210 230 Transition temp, at 50% Cry sty 39°C 35°C -5°C 17°C 71°C Transition 1 temp, at 20 f t . l b s . 25°C -5°C -45°C -10°C 51°C
10
directly applicable.
The thermal cycles occurring in the weld HAZ a r e Influenced by a number of factors
:-1. Thickness of material. 2. The welding p r o c e s s .
3. The welding conditions, current voltage and welding speed which control the heat input for unit length of weld.
4. Application of preheat or post-heat,
The results reported here a r e concerned with one set of welding
conditions; namely the submerged ar c welding of thick plate ( i , e , > 1^" thick-ness). with a heat Input of 108 kllojoules/inch, without the application of preheat.
The Iron-carbon phase diagram for equilibrium conditions and the changes In the values of the Ac. and Ac temperatures due to the rapid heating r a t e s of approximately 200 C per second occurring in the weld HAZ a r e shown in F i g . 23. The Ac^ and Ac lines were constructed from data published by Feuerstein and Smith and Albutt and Garberl2_ ^ ^ Q measured the changes
In the Ac. and Ac points of steel with varying carbon contents under different r a t e s of heating. F o r the carbon content of the steels used In this work the increases in the Ac and Ac points would be approximately 35°C and 80°C respectively.
However, the above assessment has not taken into account the Influence of the alloy content of these s t e e l s . Andrews^^ has made an intensive examination of this factor and derived the following formulae for calculating the Ac. and Ac t e m p e r a t u r e s :
-Ac = 723 - 10.7Mn - 16,9Ni + 29,1SI + 16,9Cr + 290As + 6,38W Ac = 910 - 2 0 3 / c " - 15,2Nl + 44,7Sl + 104V + 31, 5Mo + 13.IW - 30Mn
- l l C r - 20Cu + 700P + 400A1 + 120As + 400T. Applying these formulae to the steels used In this work the following values a r e obtained:-BS15 BS1501 Ac^ 715°C 723°C ^ ^ 3 821°C 831°C
The correlation of the temperature measurements made In the HAZ of BS15 mild steel bead-on-plate welds with the changes in microstructure showed that, due to the rapid heating, the Ac. point was raised by approximately 35°C to 750°C and the Ac„ point by approximately 75°C to gOO^C, These results
11
-agree with those reported by Feuerstein and Smith ,
The dependence of the Ac and Ac temperatures on heating rate Is understandable in the light of the Inability of carbon and other alloying elements to diffuse uniformly. Albutt and Garber^^ showed that for a given carbon content the distribution and shape of the carbide influenced the degree to which the Ac was increased, whereas the Ac. was insensitive to t h i s . However it Is possible that other factors, such as initial inhomogenelty and grain size may affect the values of the Ac^ and Ac t e m p e r a t u r e s .
4.1 Simulation of Thermal Cycles
Metallographic structure and hardness of the simulated regions of the weld HAZ compared favourably with the equivalent regions of the actual weld HAZ's In both m a t e r i a l s . With the BS1501 steel, although the HAZ region
chosen for observation was narrower than the similar region in BS15. the positions of the simulated thermal cycles were estimated at these points and the structure and hardness of the real and simulated regions were comparable,
A comparison between the thermal cycles measured during welding and the equivalent thermal cycles reproduced by the controlled resistance heating equipment is shown in Fig. 24, The two exanaples represent the extremes of thermal cycles used for simulation. In both c a s e s , good agreement was obtained in both the heating and cooling r a t e s , and the peak temperatures,
The measured temperature difference of 30 C between the surface and centre of the specimen during simulation was neglected".
4.2 Hardness of the Actual Weld HAZ
The hardness surveys a c r o s s the actual weld HAZ In both materials (Figs. 3 and 4) showed a scatter of approximately ± 1 0 kgms/mm^^ using an applied load of 5 kgms. This was not surprising since the structures were a mixture of carbides in a relatively soft matrix of f e r r i t e . The hardness tended to Increase on traversing from the parent plate towards the fusion boundary. This trend was expected since peak temperatures and cooling r a t e s increased In this direction and hence the proportion of the harder t r a n s -formation products. In both cases the hardest structure was associated with the region of grain coarsening.
4, 3 Banding of the Parent Plate Structures
Banding in both materials persisted to temperatures In the weld HAZ considerably above the Ac point. The problem of banding in carbon and low
14
alloy steels has been reviewed by Cairns and Charles , It Is now generally accepted that banding is due mainly to the microsegregatlon of various
elements ( e . g . phosphorus, arsenic and manganese) which diffuse slowly even at hot working t e m p e r a t u r e s . Such segregates influence the carbon distribution and form mainly during the original solidification of the ingot. The impurity and alloy elements a r e concentrated between the grains and elongate during hot
12
-working Into fibres or bands.
Although the effect of banding on Impact strength Is not completely under-stood, many high manganese steels show very marked banding and directional properties which confer high values in Impact t e s t s when the samples are cut In the direction of rolling and particularly when the notch is cut in the plane of the plate.
There Is a possible connection between cracking In the parent plate adjacent to the weld and the existence of banding in the material^'*. With fine periodic banding In mild steel the ferrite may be In appreciable constraint by the surrounding pearlite and unable to deform plastically to prevent the extension of a crack. F u r t h e r , the banding itself may be due to the s e g r e -gation of f e r r i t e formers Inherited from original solidification through the working p r o c e s s e s , such that the ferrite Is alloy embrittled,
4,4 Mechanical Properties and Microstructures in BS15 Simulated Weld HAZ Regions
A comparison of the Charpy V-notch impact properties of the four
simulated HAZ regions Is shown In F i g s , 25 and 26, and the changes in their 20 ft.lbs and 50% crystallinlty transition temperatures are shown in Fig, 27. Table 4 gives a complete summary of the mechanical testing data.
The parent material had a high transition temperature, probably because It was only slightly killed during manufacture and hence, presumably, contained a considerable amount of oxygen and nitrogen and also because it had a coarse grained s t r u c t u r e . Thermal simulation to peak temperatures of 788°C, 893°C and 1070°C produced an improvement in the material impact properties, whilst simulation to a peak temperature of 1305°C produced a severe deterioration in these properties. These changes In notch toughness could be related to the observed m i c r o s t r u c t u r e s .
(a) Thermal simulation to a peak temperature of 788°C
The peak temperature reached in this zone, which just exceeded the Ac. point under these conditions of rapid heating, had effectively limited austen-Itlsatlon to the pearlite a r e a s , with the ferrite grains remaining essentially unchanged. Since the time at which the specimen was above the Ac t e m p e r -ature was very short (approximately 8 seconds) an inhomogeneous mixture of partially transformed pearlite surrounded by high carbon austenite in a matrix of untransformed ferrite was produced. The final microstructure, shown in F i g . 5, consisted of transformed pearlite, untransformed pearlite, newly formed f e r r i t e , and untransformed f e r r i t e . The electron micrographs showed that the transformed pearlite resembled feathery upper bainite,
During the short time above the Ac. temperature there was a tendency for carbon to diffuse from the high carbon austenite to the surrounding f e r r i t e . On subsequent fairly rapid cooling this carbon could be retained In solid
13
-solution in the f e r r i t e . Work at the National Physical Laboratory showed that this had a marked effect on the Charpy V-notch properties of Iron, due to the Increased resistance to dislocation movement brought about by the increased interstitial carbon content of the f e r r i t e .
The slight improvement in notch toughness over that of the parent plate due to the thermal cycle was attributed to an increase in the overall size of the original pearlitic a r e a s , with a subsequent refinement of the original proeutectoid ferrite grains,
(b) Thermal simulation to a peak temperature of 893°C
Heating the parent material to this peak temperature, which was close to the Ac point, allowed a longer time (approximately 18 seconds) for
Ó
austenltlsatlon to develop and at a higher tenaperature than in the previous thermal cycle. The time above the Ac was sufficient for virtually complete dissolution of the pearlite to form a structure of Inhomogeneous high carbon austenite and f e r r i t e . On subsequent fairly rapid cooling the high carbon austenite transformed to fine grained ferrite and carbide as shown in Fig, 6. At high magnifications the carbide a r e a s resolved Into an upper balnlte
structure. The impact properties of this structure were considerably improved mainly due to the fine ferrite grain size,
(c) Thermal simulation to a peak temperature of 1070°C
On heating to this peak temperature, which considerably exceeded the Ac point, complete austenltlsatlon had occurred. However, the total time at which the temperature was above the Ac was too short (approximately 5 seconds) for complete homogenlsation of the austenite and for appreciable grain growth to occur. On subsequent rapid cooling through the critical temperature range, a network of proeutectoid ferrite formed at the austenite grain boundaries, and within this network a very fine mixture of ferrite and carbide with the ferrite precipitated along preferred directions. The resulting structure had a typical Widmanstatten appearance, although this was vqry fine because of the absence of marked grain growth during the austenltlsatlon part of the thermal cycle,
The Impact strength was Improved over the parent plate. In fact, the 20 ft.lbs, transition temperature was about the same as for the material simulated to a peak temperature of 788°C although the microstructure was considerably different. The Widmanstatten character of the ferrite would be expected to cause a more marked deterioration of the Impact properties than was observed, but this was probably mitigated to some extent by the fineness of the s t r u c t u r e ,
(d) Thermal simulation to a peak temperature of 1305°C
Peak temperatures In this region were greatly in excess of the Ac point. Consequently the thermal conditions were sufficient ior complete or n e a r
14
-complete homogenlsation of the austenite and for marked grain growth to occur. The subsequent cooling rate was fast enough to produce a very coarse
Widmanstatten s t r u c t u r e , as shown in Fig, 8 consisting of a network of ferrite grains outlining the prior austenite grains, adjacent to large areas of carbide and ferrite In a fine distribution. The impact properties were severely reduced due to the coarseness of the structure and the Widmanstatten f e r r i t e . The transition temperature was about 25°C higher than that of the parent plate, If the Widmanstatten ferrite were restricted from deforming by the pearlite, cleavage type of fracture could occur.
4.5 Mechanical Properties and Microstructures In BS15Q1 Simulated Weld HAZ Regions
A comparison of the Charpy V-notch impact properties of the four simulated HAZ regions is shown In F i g s , 28 and 29 and the changes in their 20 f t , l b s . and 50% crystallinlty transition temperatures in Fig, 30, The tensile properties were shown in Fig, 20 and a complete summary of the mechanical testing data in Table III.
This steel, in the "as received" condition, had a much lower transition temperature than the BS15 steel. This difference could be explained In t e r m s of the degree of deoxidation used In manufacture and the grain size of the f e r r i t e . The BS1501 was a silicon killed steel and hence the interstitial oxygen and nitrogen contents were expected to be low which would favour a low transition t e m p e r a t u r e . The ferrite grain size was considerably finer than in BS15 which would also favour a lower transition temperature,
The effect of thermalsimulation to all temperatures above the Ac. point was to increase the transition temperature of this steel. These increases varied between 7°C and 50°C depending upon the region of the HAZ being considered and the criterion used for assessing the transition temperatures, The poorest notch-toughness properties were associated with the region of grain coarsening and the partially transformed region experiencing a peak temperature of 788°C, A similar increase in Charpy V-notch impact
transition temperature in the grain coarsened region was reported by Inagakl et a l l 6 fQj. a steel of identical composition. This deterioration in impact properties was reflected in the changes occurring in the tensile properties. All the simulated weld HAZ regions showed an increase in proof s t r e s s and U , T , S , and a decrease in ductility compared with the parent material. These changes were greater n e a r e r the fusion boundary and In this respect agreed with the hardness measurements. Once again the poorest properties were associated with the region of grain coarsening,
Explanations for the changes in mechanical properties occurring in the simulated weld HAZ have already been attempted for the BS15 steel. A few differences were observed with BS1501 and the following comments supplement those already made for BS15,
(a) Thermal simulation to a peak temperature of 788 C
15
-of the ferrite grains similar to the corresponding region in BS15, -Austenlt-lsatlon was more fully developed and the resulting high carbon austenite has transformed to a feathery upper bainitic structure. Microhardness m e a s u r e -ments using the Reichert microhardness t e s t e r and a load of 27 grams, revealed the hardness of this tranformed phase was 418 kgm/mm^ compared with 275 kgm/mm for the f e r r i t e . These were only comparative values since the complete Meyer analysis, necessary to obtain absolute hardness values, was not performed.
These hardness values precluded the possibility of martensite being present and suggested the phase to be a type of upper bainite. This was supported by the tensile strength measurements. The value of 33,9 t , s , l , was in line with the tensile strength measurements of a fully upper bainitic structure of about 40 t , s , i . determined by Irvine and Pickering^',
The reduction in ductility and the deterioration in impact properties was thought to be due to the effect of this upper bainite on the structure. This was partly compensated by a reduction in the ferrite grain size. The h a r d -n e s s , o-n the other ha-nd, was comparatively little affected, probably due to the small proportion of upper bainite present in the microstructure,
(b; Thermal simulation to a peak temperature of 893 C
This structure was quite different from the equivalent one in BS15 but was similar to the previous one in BS1501, Austenitisation was more fully developed and the resulting structure had a c o a r s e r ferrite and bainite grain size than the one cycled at 788°C, The notch-toughness improved slightly due to a reduction in the carbon content of the austenite by dilution into the surrounding ferrite during the time the temperature was above the Ac
(c) Thermal simulation to a peak temperature of 1088 C
A Widmanstatten s t r u c t u r e , somewhat c o a r s e r than the equivalent specimen in BS15, was produced, Aclcular upper bainite was observed at higher manif icatlons, The transition temperature Increased slightly from that of the previous thermal cycle, but was still lower than for the sample cycled to 788°C,
(d) Thermal simulation to a peak temperature of 1347°C
A very coarse Widmanstatten structure, similar to the equivalent region in BS15 was produced, and this was associated with the highest transition temperature, greatest strength, and lowest ductility in the weld HAZ, This was also reflected in the marked reduction in the U , T , S , / P r o o f s t r e s s ratio, as shown in Table III, More aclcular upper bainite than in the previous thermal cycle was observed at higher magnifications,
4.6 The Effect of Restraint on the Simulated HAZ Properties of BS15 During welding the HAZ undergoes an extremely complex strain cycle and the effects of this on the structure and properties a r e uncertain. Several
16
-s t r e -s -s and con-sequent pla-stic deformation on the i-sothermal tran-sformation of austenite to ferrite and pearlite, bainite, and martensite. They found that the applied s t r e s s shortened the beginning and ending times of transformation and accelerated the rate of transformation,
24
Smallman claimed that even elastic s t r e s s e s , when applied above the M temperature and maintained during cooling, can affect the transformation
6
behaviour. Uniaxial compression or tensile s t r e s s e s raise the M temperature,
s
while hydrostatic s t r e s s e s lower the M temperature.
Hence it is possible that the temporary dynamic thermal s t r e s s e s and strains produced in the HAZ during welding could significantly affect the resultant microstructure and its mechanical properties. This is particularly so in the region of a defect where the s t r e s s may be raised by as much as three t i m e s . However, techniques to simulate HAZ structures containing defects were not employed,
Because of the absence of reliable data on the strain cycles occurring in the HAZ, this problem has been tackled In a qualitative manner by applying a restraint during thermal simulation. No detectable changes in microstructure or mechanical properties was observed C. F i g s , 15 and 16,
This evidence is not conclusive since the nature and magnitude of the strains imposed in this approach differ from those occurring during welding.
Stress may also play an Important part in the diffusion of hydrogen from the weld bead into the surrounding HAZ, The effect of hydrogen as an embrittling medium is well known, but the mechanism of embrittlement is still uncertain. This aspect is beyond the scope of the present work,
4,7 The Effect of Post-Weld Heat Treatment
A subsequent heat treatment to 650 C±50 C for 30 minutes on the fracture toughness properties of simulated weld HAZ structures in BS15 led to very little change in transition t e m p e r a t u r e s . Unfortunately, the region of greatest interest, l , e . the region of grain coarsening, was not investigated due to shortage of material,
These r e s u l t s indicate that this post-weld heat treatment is not beneficial for mild steel weldments, except possibly in the case of very thick sections where residual s t r e s s e s may be high.
5, Conclusions
1, The non-equilibrium conditions attending a submerged arc bead-on-plate weld In mild steel at a heat input of 108 kllojoules/inch resulted In an increase of approximately 35°C In the Ac temperature and of approximately 75°C in the Ac temperature,
17
-2. The weld HAZ of mild steel could be divided into four distinct regions, These were: (a) the region of grain coarsening, (b) the region of grain refine-ment, (c) the region of partial transformation and (d) the region of spheroidis-ation,
3. The peak hardness occurred In the region of grain coarsening and there was a general decrease in hardness with increasing distance from the fusion boundary,
4. Metallographic examination and hardness measurements showed that the simulated HAZ regions were comparable with the actual HAZ regions In both m a t e r i a l s ,
5. An Imposed restraint during thermal simulation produced plastic deform-ations of 2-4% but had no detectable influence on the final microstructure and mechanical properties of BS15.
6. BS15 parent material simulated to peak temperatures of 788°C, 893 C, and 1070°C showed an improvement in Charpy V-notch toughness as compared with the parent plate. This was attributed to a decrease in the ferrite grain size. The material simulated to a peak temperature of 1305°C had consid-erably lower impact strength than the parent material due to a very coarse grained structure consisting of large a r e a s of fine carbide surrounded by coarse Widmanstatten f e r r i t e .
7. The BS1501 parent material simulated to all four peak temperatures above the Ac^ point showed a decrease in notch toughness. This was most severe in the grain coarsened region and the partially transformed region simulated to a peak temperature of 788°C, The embrittlement in the partially transformed region was attributed to the formation of feathery upper bainite in the former pearlitic a r e a s . The embrittlement in the grain coarsened
region was attributed to the formation of a very coarse Widmanstatten structure of proeutectoid ferrite and aclcular upper balnlte,
8. The tensile properties of the BS1501 parent material changed continuously with increasing peak temperature of simulation. The highest values of proof s t r e s s , tensile strength, and hardness, and the lowest values of ductility, occurred in the region of grain coarsening,
18
-6, B i b l i o g r a p h y
1. N i p p e s . E , F , . Welding J , . 38 (1). 1959, I s - 1 3 s ,
2 . Wilson, J . E . B . , D . A , E , T h e s i s , C o l l e g e of A e r o n a u t i c s , C r a n f i e l d , 1964.
3 . A p b l e t t , W , R , , et a l . Welding J , , 31_ (11), 1952, 421s - 4 3 0 s ,
4 . S e f e r i a n , D . , " T h e M e t a l l u r g y of Welding", Chapman and H a l l , 1962. 5 . L a n c a s t e r , J , F . , " T h e M e t a l l u r g y of Welding, B r a z i n g and S o l d e r i n g " ,
G e o r g e Allen and Unwin L t d . , 1965,
6. C o w a r d , M . D , and A p p s , R . L , , CoA Note M a t , N o . 1 3 , College of A e r o n a u t i c s , C r a n f i e l d , 1967.
7 . G e o r g e , M , J , and Clifton, T , E . , CoA N o t e , College of A e r o n a u t i c s , C r a n f i e l d , To be published.
8. C a l v o , F . A . et a l . Studies of t h e Welding M e t a l l u r g y of Steel, B . W . R . A . , 1 9 6 3 . 9. N i p p e s , E . F . and S a v a g e , W , F , , Welding J , , 28^(11) 1949, 534s - 5 4 6 s . 1 0 . B r o w n , L , J , , M , S c T h e s i s , College of A e r o n a u t i c s , C r a n f i e l d , 1967. 1 1 . F e u e r s t e i n , W , J . and S m i t h , W , K , , T r a n s , A . S , M , 46 1954, 1270. 12. Albutt, K , J , and G a r b e r , S. , J , I , S , I . 204 (12) 1966, 1217, 1 3 . A n d r e w s , K , W . , J , I , S , I , 203^(7) 1965, 7 2 1 ,
1 4 . C a i r n s , R . L . and C h a r l e s , J , A , , I r o n and S^eel, Nov, 1966, 5 1 1 - 5 1 5 . 1 5 . A l l e n . N , P , , J , I , S , I , , 174, 1953, 108. 1 6 . I n a g a k l , M. et a l . T r a n s , N a t . R e s . I n s t , M e t , ( J a p a n ) 6 (6), 1964, 7 3 - 8 3 . 1 7 . I r v i n e , K . J , and P i c k e r i n g , F , B , , J , 1 , S , I . 2 0 1 , (6) 1963. 518. 1 8 . B i r k s , L . B , , T r a n s . A , I , M . E . , 1955, 179, 1 9 . B h a t t a o h a r y y a , S, and K e h l , C . L . , T r a n s . A , S , M , 47^, 1955, 3 5 1 , 2 0 . K e h l , C , L , and B h a t t a o h a r y y a , T r a n s . A . S , M , 4 8 , 1958. 234, 2 1 . P o r t e r , L , F . and R o s e n t h a l , P . C , Acta M e t , 7^, 1959, 504, 2 2 . K a n a z a w a , S. , T r a n s , J a p . I . M , 4 , 1 9 6 3 , 1 9 5 , 2 3 . W e l l s , M . G . H , and West, D , R . F . , J , 1 , S , 1 , 200 (9), 1962, 710, 2 4 . S m a l l m a n , R , E , , " M o d e r n P h y s i c a l M e t a l l u r g y " , B u t t e r w o r t h s , 1965.
0. 5 1. O 1. 5 2 . O
F I G . 1 THE H E A T A F F E C T E D ZONE MICROSTRUCTURES ASSOCIATED WITH
DISTANCE F R O M FUSION BOUNDARY, MM, 2. 5 3. 0 3. 5 4. 0 4. 5 5. 0 X 70 (B) (A) ' ( d ) X 700 A B E A D - O N - P L A T E WELD IN BS15 MILD S T E E L
2 . O 2 . 5
DISTANCE F R O M FUSION BOUNDARY. MM. 3 . 0 3 . 5 4 . 0 X 7 Ö SPHEROID ISATION REGION
PARENT PLATE
KTOO A B E A D - O N - P L A T E WELD IN B S 1 5 0 1 MILD S T E E L240 230 220 > 210 200 190
1
[ 1 ^--X J V \ \ I\ \
\
1 T/^^-^
r J • - AVERAGE OF k ^ 10 READINGS/ \
T
• >--4 • « Tv L__. __ _J r . . 1.'f-.
t-r
[
J
.1
-0 1 2 3 4 5 6 7 mmDISTANCE FROM FUSION BOUNDARY
FIG. 3 MILD STEEL HEAT AFFECTED ZONE HARDNESS SURVEY BS 15. 230 1 220 210 aoo in a; X IQfl lao 170 )} r l \ _ V \ 1 ( \ V i I i 0 AVERAGE OF 5 REAQNGS 1 1 — 1 I ') \ \ i 0 < ~~1 L • - - - ( I N » 1 2 3 i 5 6 7mm. DISTANCE FROM FUSION BOUNCWRY
rw\
O P T I C A L MICROGRAPH X 700 (A) - t r a n s f o r m e d p e a r l i t e (B) - u n t r a n s f o r m e d p e a r l i t e (C) - newly f o r m e d f e r r i t e (D) - u n t r a n s f o r m e d f e r r i t e (A) E L E C T R O N MICROGRAPH X 12, 000 (A) (B) E L E C T R O N MICROGRAPH X 22. 000 HV5 2 0 0 - 10F I G . 5 PHOTOMICROGRAPHS O F THE S T R U C T U R E S PRODUCED IN B S 1 5 MILD S T E E L SIMULATED TO A PEAK T E M P E R A T U R E
(B)
O P T I C A L MICROGRAPH X 700 (A) - ferrite (B) - u p p e r bainite E L E C T R O N MICROGRAPH X 4 , 0 0 0 E L E C T R O N MICROGRAPH X 20, 000 HV5 205 - 10
F I G . 6 PHOTOMICROGRAPHS O F THE STRUCTURES PRODUCED IN B S 1 5 MILD S T E E L SIMULATED TO A P E A K T E M P E R A T U R E OF 893"C
(A) - p r o e u t e c t o i d f e r r i t e ; (B) - fine carbide and f e r r i t e
F I G . 7 PHOTOMICROGRAPH OF BS 15 MILD S T E E L SIMULATED TO A P E A K T E M P E R A T U R E O F 1, 070 C
(A)
HV5 230"^ 10
(A) - p r o e u t e c t o i d f e r r i t e ; (B) fine c a r b i d e and f e r r i t e
F I G . 8 PHOTOMICROGRAPH O F B S I 5 MILD S T E E L SIMULATED TO A P E A K T E M P E R A T U R E O F 1 , 3 0 5 ° C
HV5 199 - 10
x 7 0 0
(A) f e r r i t e , (B) p e a r l i t e
OPTICAL MICROGRAPH x 7 0 0
(A) (B)
(A) - ferrite
(B) - feathery upper bainite
ELECTRON MICROGRAPH X 3 2 , 0 0 0 HV 10 1 7 0 - 10 ELECTRON MICROGRAPH X 1 4 0 0 0
FIG, 10 PHOTOMICROGRAPHS OF S T R U C T U R E S PRODUCED IN BS 1501 MILD STEEL SIMULATED TO A P E A K T E M P E R A T U R E O F 788°C
MICROGRAPH X 700 E L E C T R O N MICROGRAPH X 18, 000 (A) - f e r r i t e (B) - u p p e r b a i n i t e E L E C T R O N MICROGRAPH X 2 3 , 0 0 0 HV 10 . 1 9 0 ' t 10 F I G , 11 PHOTOMICROGRAPHS O F S T R U C T U R E S PRODUCED IN B S 1 5 0 1 MILD S T E E L SIMULATED TO A PEAK
O P T I C A L MICROGRAPH X 700 J. , .«..'I t «v;- ' • CB) E L E C T R O N MICROGRAPH x l 5 0 0 (A) - p r o e u t e c t o i d f e r r i t e (B) - u p p e r b a i n i t e E L E C T R O N MICROGRAPH X 1400 HV 10 207"*^ 10
F I G . 12 PHOTOMICROGRAPHS O F STRUCTURES PRODUCED IN B S 1 5 0 1 MILD S T E E L SIMULATED TO A PEAK T E M P E R A T U R E O F l , 0 8 8 " c
O P T I C A L MICROGRAPH X 700 (A) (B) (B)'
T7
r-'f\
c
• . . « / • '.^-v» •^;
^^ ''y
N V (A) - p r o e u t e c t o i d f e r r i t e (B) - u p p e r b a i n i t e?5^
E L E C T R O N M I C R O G R A P H X 7, 000 E L E C T R O N MICROGRAPH X 26,000 HV 10 231 ""- 10:;:v^^§i'
' ^ ' (B)FIG. 13 PHOTOMICROGRAPHS O F STRUCTURES IN B S 1 5 0 1 Mli.D S T E E L SIMULATED TO A PEAK T E M P E R A T U R E O F 1, 347°C
HV 10 1 38 ^ 10
X 7 0 0
(A) - f e r r i t e . (B) - c a r b i d e .
FIG. 14 PHOTOMICROGRAPH OF BS 1501 MILD S T E E L P A R E N T P L A T E
-100 - 6 0 -20 20 60 100 -100 - 6 0 -20 UO
TESTING T E M P E R A T U R E , °C
H G , 1 5 . CHARPY V-NOTCH IMPACT (FT-LBS) TEMPERATURE DATA FOR MILD STEEL SIMULATED SPECIMENS BS 15.
20 40 100 HO TESTING TEMPERATURE. °C
FI6.16 CHARPY V-NOTCH IMPACT (PERCENTAGE CRYSTALLINITY)-TEMPERATURE DATA FOR MILD STEEL SIMULATED SPECIMENS BS 15.
-100 -80 -20 20 60 100 -60 -20 20 60 100 140 TESTING TEMPERATURE , °C.
FI6.17 CHARPY V-NOTCH 1MR4CT (FT-LBS) TEMPERATURE DATA FOR B.S.1501 MILD STEEL
SIMULATED SPECIMENS.
• 1 0 0 - 6 0 -20 20 ' ' 60 80 -60 -20 20 60 100 W
TESTING TEMPERATURE, °C ^ ^ ^ FIG 18. CHARPY V-NOTCH IMPACT ('/.CRYSTALLINITY)-TEMPERATURE DATA FOR B.S.1501 MILD STEEL
UI z o.
o
ZERO RESTRAINT
FIG 19 MEASUREMENT OF THE DISPLACEMENT CF COPPER BLOCKS DURING THERMAL CYCLING TO 1070'C UNDER CONDITIONS OF ZERO AND FULL RESTRAINT.
0 1 2 3 4 5 6 7 DISTANCE FROM FUSION BOUNDARY(mm)
FIG20. TENSILE PROPERTIES CF BS 1501 MILD STEEL SIMULATED SPECIMENS.
i
s
- 1 UJ
788° C THERMAL
os cyclsd
_ _ c y c l e d . p w h t
lOTO'C THERMAL CVCLE
jrjx'
1305 "C THERMAL CYCLE as cycled
I
-loü ^6ö ^2Ö M w iSo T45 m • ^53 ^20 '. S T Ö iSi5 lïC i W
TEMPERATURE, °C
FIQ 21. EFFECT OF POST V ^ L Ü HEAT TREATMENT (30 MINS. AT 65CPC - 50°C) ON WELD KA.Z. PROPERTIES CF BS 15 MILD STEEL.
-100 -60 20 20 60 100 " 180 -60 -20 20 60 «O
TEMPERATURE °C , ^
FIG22 EFFECT OF POST V^LD HEAT TREATMENT (30 MINS AT 650°C150°C) ON WELD HAZ PROPERTIES OF B.S.15 MILD STEEL.
RESULTS of FEUERSTEIN and SMITH RESULTS of ALBUTT end GARSER" RESULTS of present work.
.10
893C
0-8 1-0 CARBON CONTENT,* MILD STEEL and
QT35 CARBON EQUIVALENT
FIG. 23 THE I RON-CARBON PHASE DIAGRAM WITH THE Ac, and Acj LINES CONSTRUCTED FOR A
HEATING RATE of 200*C PER SECOND.
^
\f
11
1 1 1 'ij
11 1 A 1 A 1 // /' 1 /' '1 ['1
t\f
' \ \V
\ AS SIMU V \ vs.1 1 1 1
*<EASURED THERMAL CYCLE
\
v\
\ ^ ^ . •V 5 » ^ ^ ^*'*'^,^^— • I
1347°C THERMAL 1 CYCLE1
;">^-^-7e8°C THERMAL 10 20 30 40 50 60 70 80 100 TIME/SECSl O O i
-60 80 100 120 TEST TEMPERATURE °C
F I G . 2 5 COMPARISON OF CHARPY V-NOTCH IMPACT (PERCENTAGE CRYSTALLINITY)
-TEMPERATURE CURVES FOR MILD STEEL SPECIMENS BS15
1305°C
80 100 120 140 C
TESTING TEMPERATURE
F(G26 COMPARISON OF CHARPY VNOTCH IMPACT ( F T L B S )
-TEMPERATURE CURVES FOR MILD STEEL SPECIMENS BS15.
IMPACT STRENGTH (FT-LBS) TRANSITION T E M P E R A T U R E , °C 2 Ki 2 < o o 5m 5^ mm —i 3 m >
pa
" 1 — t/>m > ^ 7) 8S! ^ / / / / / / / / / / 1 / 1J
DUCTILIT Y TRANSITIO N TEMPERATUR E A T 2 0 FT-LB S LEVE L1
FRACTUR E TRANSITIO N TEMPERATUR E A T 50 % rBV<!TALLINIT Y N . XFRACTURE APPEARANCE (PEHCBnA6E GRVSTALLWITY)
o g S 8 8 g 8 S ^ O 2;
