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THE COLLEGE OF AERONAUTICS
CRANFIELD
THE WELD HEAT A F F E C T E D ZONE STRUCTURE
AND PROPERTIES O F QT 35 STEEL
by
THE COLLEGE OF AERONAUTICS
CRANFIELD
T h e Weid H e a t Affected Zone S t r u c t u r e and P r o p e r t i e s of QT 35 S t e e l
by
M. D. C o w a r d , M. P h i l . . B. Sc. , A. R. S. M. , E. Smith, P h . D. , B. Sc. , A. I. M.
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 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 w a s a c h i e v e d 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 c a p a b l e of cycling s p e c i m e n s t h r o u g h 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 1^ inch thick p l a t e u s i n g a h e a t input of 108 k i l o j o u l e s p e r inch.
M e t a l l u r g i c a l e x a m i n a t i o n and h a r d n e s s m e a s u r e m e n t s showed that the s i m u l a t e d r e g i o n s w e 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 HAZ. T h e C h a r p y V - n o t c h i m p a c t 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 w e r e m e a s u r e d and c o r r e l a t e d with 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 o p t i c a l 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 showed t h a t 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 r e d in all r e g i o n s of the v i s i b l e HAZ. T h i s w a s m o s t s e v e r e in the g r a i n c o a r s e n e d r e g i o n and w a s a s s o c i a t e d with the f o r m a t i o n of a c o a r s e g r a i n e d u p p e r b a i n i t i c m i c r o s t r u c t u r e .
A r e s t r a i n t t e c h n i q u e giving p l a s t i c d e f o r m a t i o n s of 2-4%, applied d u r i n g t h e r m a l cycling w a s shown to have no significant effect on the r e s u l t i n g m i c r o -s t r u c t u r e and p r o p e r t i e -s .
A p o s t - w e l d h e a t t r e a t m e n t at 150°C for 4 h o u r s on s p e c i m e n s s i m u l a t i n g the p a r t i a l l y t r a n s f o r m e d and g r a i n refined r e g i o n s was shown to have a slight b e n e f i c i a l effect on the f r a c t u r e t o u g h n e s s p r o p e r t i e s at low t e m p e r a t u r e s .
F i n a l l y , s u g g e s t i o n s a r e m a d e for m i n i m i s i n g the e m b r i t t l e m e n t o c c u r r i n g in the HAZ of high h e a t input welds in QT 35.
Introduction 1 Experimental 2 2. 1 Materials 2 2. 2 Experimental Procedure 2
2. 2. 1 Preparation and Examination of Weld 2
2. 2. 2 Simulated HAZ Structures 3 2. 2. 3 Post-Weld Heat Treatment 3
Results 3 3. 1 Metallography and Hardness of the Weld HAZ 3
3. 2 Metallography of the Simulated HAZ Structures 4 3. 3 Mechanical P r o p e r t i e s of the Simulated HAZ Structures 5
3. 4 Effects of the Post-Weld Heat Treatment 5
Discussion 5 4. 1 Metallography of the Weld HAZ 5
4. 2 Hardness of the Actual HAZ 6 4. 3 Mechanical P r o p e r t i e s of the Simulated HAZ Structures 6
4. 4 The Effect of Restraint on the Simulated HAZ
Structures 7 4. 5 The Effect of a Post-Weld Heat Treatment 7
4. 6 The Importance of the HAZ 7
Conclusions 8 Bibliography 9 Figures
1. Introduction
QT 35 is a medium strength, low carbon, low alloy (Mn-Ni-Cr-Mo-V). quenched and tempered steel which was developed for use in nuclear submarine hulls and highly s t r e s s e d parts of surface ships, where the ability to deform in a ductile manner at low temperatures is desirable. Since the majority of its applications involve welding fabrication, it is important to know the effects of welding on the parent plate properties.
The thermal cycles occurring in the vicinity of a weld alter the m i c r o -structure and mechanical properties of the steel significantly. The effect on notch-toughness is of particular importance since for QT 35 this determines the capacity of the steel to withstand impulsive loading at moderately low temperatures. In addition this tjrpe of steel is susceptible to HAZ cold c r a c -king when the following three conditions a r e satisfied
simultaneously^"3:-1. The presence of a susceptible microstructure. 2. The presence of hydrogen.
3. The presence of s t r e s s .
The practical methods of avoiding this type of cracking are based on the control or elimination of one or more of these three factors. The present work focuses attention on the microstructure and its effect on mechanical properties.
The original microstructure contains tempered bainite and/or martensite which give good low temperature properties. Nippes ^' ^, Arnold and Savage and Owczarski "^, however, have shown that in this class of steel a distinct impairment of notch toughness occurs in the HAZ of high heat input welds due to the effects of slow r a t e s of cooling on the microstructure. The greatest damage was observed in the grain coarsened region. A niunber of explanations have been advanced. Savage and Owczarski ''. Nisbett and Kelly °, and
Oldridge " attributed this embrittlement in the grain coarsened region to the formation of upper bainite. Winn and Wortley 1^ observed a coarsening of the ferrite laths with carbide precipitated within and at the interfaces of the laths. Bentley and Smith ^^ observed a coarse bainite with more free carbides and less martensite. Savage and Owczarski ' also observed a severe embrittle-ment in the intercritical region due to the formation of a mixture of ferrite and martensite.
The aim of the present work was to examine the range of microstructures produced in the HAZ of a submerged arc bead-on-plate weld in 1^ inch thick QT 35 plate using a high heat input of 108 kilojoules per inch. The practical difficulties associated with the accurate positioning of test specimens due to the wide variety of structures occurring over a very narrow region were overcome by using a simulation technique, described in a previous report ^^, to produce specimens for the determination of the HAZ properties. Details of the thermal cycles used were taken from the work of Coward and Apps '••^.
The results showed that the visible weld HAZ was made up of three distinct regions. These were the region of grain coarsening, the region of
grain refinement, and the region of partial transformation. Metallurgical examination and hardness measurements showed that the simulated structures were comparable with those found in the weld HAZ. All the sintiulated HAZ s t r u c t u r e s had reduced notch-toughness properties compared to the parent plate. The greatest embrittlement was observed in the grain coarsened region and this was associated with the presence of upper bainite in the microstructure.
The effects of a heavy r e s t r a i n t applied during the simulation of the thermal cycles was also investigated. The applicability of this treatment to the welding situation was discussed in a previous report 1^. No measurable effects on mechanical properties and m i c r o s t r u c t u r e s were observed.
The effects of a post-weld heat treatment at 150°C for 4 hours on specimens simulated through weld thermal cycles with peak temperatures of 788°C, 8930C and 1070°C were studied. The results showed a slight improve-ment in notch-toughness at low temperatures although in some cases this was accompanied by a slight reduction in notch-toughness at high temperatures. Unfortunately, the grain coarsened region was not investigated due to a short-age of material.
2. Experimental 2. 1 Materials
The m a t e r i a l was supplied as l | inch thick plate. The chemical analysis is shown in Table 1. The miaterial specification gives a nominal hardness of 230-250 HV and a 0.2% proof s t r e s s of 36-44 t. s. i. In addition, a Charpy V-notch toughness of 40 ft. lbs. minimum at -40°C. with a fracture appearance of not g r e a t e r than 75% crystallinity. is specified.
C 0. 13 Mn 0.97 Si 0. 16 S 0. 02 P 0,022 Ni 1. 06 C r 0.77 Mo 0.39 V 0.06 TABLE 1 Chemical analysis of QT 35 plate 2. 2 Experimental Procedure
The experimental procedure was similar to that described in a previous report ^^. A summary only is given here.
2. 2. 1 Preparation and Examination of Weld
A bead-on-plate weld was made using a submerged a r c welding unit and a heat input of 108 kilojoules per inch. A transverse section through the weld HAZ was examined metallographically and hardness determinations made at 0. 5 mm. intervals starting from the fusion boundary using a Zwick hardness tester and a load of 5 kilograms.
2. 2. 2 Simulated HAZ Structures
Structures representing four different parts of the visible weld HAZ were reproduced in specimens 1. 0 cm. square and 7. 5 cm. long using the simulation technique. In most tests no r e s t r a i n t was applied to the specimen during sinnulation, but in some tests a r e s t r a i n t was applied to prevent
expansion and contraction. The peak temperatures associated with these structures were 788°C, 893°C. 1070°C and 1347°C. The structures were examined metallographically and compared with the structures observed in the actual weld HAZ. The mechanical properties of these structures were determined by means of hardness and Charpy V-notch impact tests and the results compared with the parent plate properties.
2. 2. 3 Post-Weld Heat Treatment
HAZ structures produced by simulation to peak temperatures of 788°C, 893°C and 1070°C were given a post-weld heat treatment at 150°C for 4 hours in a hot-air circulation furnace. Charpy V-notch impact tests were then c a r r i e d out and the results compared with those obtained without a post-weld heat treatment.
3. Results
3. 1 Metallography and Hardness of the Weld HA.Z
The parent plate microstructure was fine grained and consisted of a distribution of alloy carbides in a ferrite matrix resulting from a tempering of a mixture of upper and lower bainite and possibly some martensite. This structure is shown in Fig. 1.
The range of m i c r o s t r u c t u r e s produced by the bead-on-plate weld is shown in Fig. 2. The HAZ was divided into three distinct regions and these were identified in t e r m s of their distance from the weld fusion boundary as follows
:-(a) The region of grain coarsening
This region extended from 0 to 0. 8 mm. from the fusion boundary and experienced peak temperatures in excess of about 1100°C. Since the max-imum temperatures were well into the austenite region, austenitic grain growth took place, the effect being less marked with increasing distance from the fusion boundary. On subsequent rapid cooling an upper bainite structure formed, as shown in Fig. 2A.
(b) The region of grain refinement
This region extended from 0. 8 mm. to 2. 0 mm. from the fusion boundary and experienced maximum temperatures in the approximate range of 1100°C to 900°C. Complete austenitization occurred but insufficient time was available for homogenization of the structure. Subsequent cooling prod-uced a structure consisting of upper bainite needles nucleated at the prior
austenite grain boundaries together with a r e a s of fine carbide precipitated in a lower bainite transformation product. This is shown in Fig. 2B.
(c) The region of partial transformation
This region extended from 2. 0 mm. to 3. 1 mm. from the fusion
boundary and experienced maximum temperatures in the approximate range of 900^0 to 750°C, the temperature limits of Ac and Ac . P a r t i a l
austenitiz-X o
ation occurred producing regions of high carbon austenite which transforined on cooling to a lower bainitic product, consisting of very fine carbides p r e -cipitated within the bainitic ferrite. This is shown in Fig. 2C.
Beyond this region structural changes were not observed by optical microscopy and, therefore, limited the region defined as the weld HAZ. No region of spheroidisation was formed since the carbides in the parent plate were already spherical.
The variations in hardness a c r o s s the HAZ are shown in Fig. 3. A considerable increase in hardness occurred in all regions of the HAZ and, in general, the hardness increased with decreasing distances from the fusion boundary. Maximum hardness was located in the grain coarsened region.
3. 2 Metallography of the Simulated HAZ Structures (a) Simulated to a peak teniperature of 788°C
Thermal cycling to this peak temperature produced a microstructure consisting of newly formed lower bainite in the a r e a s which underwent austenitization and partially dissolved tempered alloy carbides in the untrans-formed matrix. This is shown in Fig. 4.
(b) Simulated to a peak temperature of 893°C
Thermal cycling to this peak temperature produced a microstructure consisting of lower bainite in which the carbides were c o a r s e r than those found in the previous sample and with the adjacent regions of ferrite almost devoid of any fine tempered alloy carbides. This is shown in Fig. 5.
(c) Simulated to a peak temperature of 1070 C
Thermal cycling to this peak temperature produced a microstructure consisting of a r e a s of upper bainite needles nucleated at the prior austenitic grain boundaries together with a r e a s of fine carbides precipitated in a lower bainitic transformation product. This is shown in Fig. 6.
(d) Simulated to a peak temperature of 1347'-'C
Thermal cycling to this peak temperature produced a uniform grain coarsened microstructure consisting of upper bainite. This is shown in Fig. 7.
3. 3 Mechanical P r o p e r t i e s of thie Simulated HAZ Structures
The hardnesses of the different .microstructures are recorded with the micrographs in Figs. 4-7. The Charpy V-notch impact results are shown in Figs. 8 and 9. The results obtained with a restraint applied during simulation a r e also included. The amount of plastic deformation introduced by the
restraining mechanism was shown to be 2% for a peak temperature of 788°C, 3% for 893°C and 4% for 1070°C 1'*.
3. 4 Effects of a Post-Weld Heat Treatment
The results of the post-weld heat treatment are shown in Figs. 10 and
11.
4. Discussion
The m i c r o s t r u c t u r a l changes observed in the weld HAZ of steels can be related to the effective Ac and Ac temperatures. The equilibrium values
i. O
of Ac and Ac calculated from A n d r e w s ' ^ formulae for the effect of
chem-J. O
ical composition a r e 712°C and 820°C respectively. Previous work •'•^ showed that these temperatures were raised by approximately 35°C and 75°C respect-ively under the conditions of rapid heating associated with the weld thermal cycles used here. Thus, the approximate values of the effective Ac and Ac„ temperatures for QT 35 a r e 747°C and 895°C respectively. Calculation has. of course, neglected the effects of shape and distribution of carbides, inhom-ogeneity and grain size which may influence these temperatures, particularly the Ac . Within these limitations, these values a r e now used to establish the relationship between the weld thermal cycle and the microstructural changes.
4. 1 Metallography of the Weld HAZ
The first observable changes in microstructure occurred when the peak temperature exceeded the effective Ac . Between this temperature and the effective Ac austenite nucleated at temper carbides located at the p r i o r austenite grain boundaries. The extent of the austenitization increased with increasing temperature until at the effective Ac austenitization was complete. Because of the very short time at which the temperature was above the
effective Ac the austenite retained its high carbon content. On subsequent cooling the austenite transformed to a lower bainitic product with very fine carbides precipitated within the bainitic ferrite. This structure is shown in Fig. 2C. In the untransformed ferrite a r e a s the fine alloy carbides had been partially taken into solution and retained there on subsequent cooling, thus strengthening the ferrite.
The structure of this region is s i m i l a r to one reported by Irvine and Pickering ^^ for a 0.61% carbon bainitic steel. The effect of carbon content
on the transformation c h a r a c t e r i s t i c s of a low alloy steel of the QT 35 type has not yet been established. In general, however, the effect of carbon and most adloying elements is to delay the initiation of transformation and to d e c r e a s e the rate at which the reaction proceeds. Hence it is probable that as the carbon content of the austenite is increased the final transformation products will change from upper to lower bainite.
As the peak t e m p e r a t u r e s approached the effective Ac the austenite contained smaller amounts of carbon and on subsequent cooling the carbides precipitated inside the bainitic ferrite became c o a r s e r and more widely spaced. Dissolution of the fine alloy carbides in the untransformed ferrite was ntiore advanced so that these regions became nearly devoid of carbide particles.
On heating above the effective Ac a completely austenitic structure formed, but below about 1100°C complete homogenisation of the austenite did not take place due to the short period of time available for carbon diffusion at these temperatures. On subsequent cooling the low carbon a r e a s t r a n s -formed to upper bainite and the high carbon a r e a s to lower bainite. Fig. 6 typifies this kind of structure.
At temperatures above about 1100°C sufficient time was available for virtually complete homogenisation of the austenite and for grain growth to occur. This became more advanced on approaching the fusion boundary. Subsequent cooling of the comparatively low carbon austenite produced an almost completely upper bainitic structure, as shown in Figs. 2A and 7.
4. 2 Hardness of the Actual HAZ
The hardness survey a c r o s s the HAZ (Fig. 3) showed a considerable increase in hardness in all regions of the HAZ. The general trend was for the hardness to increase on traversing towards the fusion boundary. In the intercritical region this was due to a strengthening of the ferrite matrix as described above and the formation of lower bainite from high carbon austenite. Above the Ac the hardness increased further due to the formation of upper
Ó
bainite in the m i c r o s t r u c t u r e . The greatest hardness was associated with a fully upper bainitic structure.
4. 3 Mechanical P r o p e r t i e s of Simulated HAZ Structures
A comparison of the Charpy V-notch impact properties of the four simulated HAZ regions is shown in Figs. 12 and 13. and the changes in their 20 ft. lbs and 50% crystallinity transition temperatures in Fig. 14. The parent plate had a very low transition temperature, typical of a tempered martensitic and/or bainitic structure. Although this steel was claimed to be a quenched and tempered martensitic steel. M o r r i s 1''. who derived the isothermal t r a n s -formation diagram for this steel, (Fig. 15), showed that the steel was not fully hardenable even in spray-quenched l|- inch thick plate. The material used for this work was cut from the centre of 1^ inch thick plate and, t h e r e -fore, the structure of tempered bainites with a possibility of some tempered martensite was in accordance with Morris.
T h i s was s u b s t a n t i a t e d by t h e continuous cooling t r a n s f o r m a t i o n d i a g r a m , F i g . 16. which shows that a cooling r a t e of m o r e than 75°C p e r s e c o n d i s n e c e s s a r y d u r i n g cooling f r o m 1000°C to 200°C to e n s u r e t h a t a c o m p l e t e l y m a r t e n s i t i c s t r u c t u r e i s f o r m e d . T h i s d i a g r a m w a s p r o d u c e d by the B r i t i s h Welding R e s e a r c h A s s o c i a t i o n ^^ by m e a n s of high s p e e d d i l a t o m e t r y from a s p e c i m e n of the QT 35 u s e d in t h i s w o r k .
T h e effect of t h e r m a l s i m u l a t i o n to peak t e m p e r a t u r e s of 788°C. 893°C. 1070 C and 1347°C w a s to p r o d u c e a m a r k e d d e t e r i o r a t i o n in i m p a c t p r o p e r -t i e s . which i n c r e a s e d in s e v e r i -t y wi-th i n c r e a s i n g peak -t e m p e r a -t u r e . The r e a s o n s for t h e s e c h a n g e s in n o t c h - t o u g h n e s s a r e d i s c u s s e d below in t e r m s of t h e i r m i c r o s t r u c t u r e s .
(a) T h e r m a l s i m u l a t i o n to a peak t e m p e r a t u r e of 788°C
T h i s s t r u c t u r e , shown in F i g . 4. r e p r e s e n t s the low t e m p e r a t u r e p a r t of 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 in which the peak t e m p e r a t u r e exceeded the effective Ac . T h e i n c r e a s e d t r a n s i t i o n t e m p e r a t u r e w a s a t t r i b u t e d to the g r e a t e r e a s e of c r a c k n u c l e a t i o n and p r o p a g a t i o n in t h i s m i c r o s t r u c t u r e . The p a r t i a l d i s s o l u t i o n of the c a r b i d e s in the f e r r i t e led to an i n c r e a s e in the
f e r r i t e m e a n f r e e path and a c o n s e q u e n t i n c r e a s e in the s i z e of d i s l o c a t i o n p i l e - u p s . The s t r e n g t h e n i n g of the f e r r i t e m a t r i x , which w a s r e f l e c t e d in the h a r d n e s s v a l u e s , inhibited p l a s t i c flow f r o m o c c u r r i n g to r e l i e v e s t r e s s c o n -c e n t r a t i o n s at the t i p s of d i s l o -c a t i o n p i l e - u p s and so m a d e f r a -c t u r e initiation conditions e a s i e r . A l s o , s i n c e only c e r t a i n a r e a s contained a high d e n s i t y of c a r b i d e s which could act a s c r a c k a r r e s t e r s , c r a c k p r o p a g a t i o n would be e a s i e r .
(b) T h e r m a l s i m u l a t i o n to a peak t e m p e r a t u r e of 893°C
T h i s s t r u c t u r e , shown in Fig. 5, r e p r e s e n t s 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 in which the peak t e m p e r a t u r e w a s c l o s e to the effective Ac . The i n c r e a s e d t r a n s i t i o n t e m p e r a t u r e w a s a t t r i b u t e d to a d e c r e a s e in the r e s i s t a n c e to c r a c k initiation due to an i n c r e a s e in the f e r r i t e g r a i n s i z e and a s t r e n g t h -ening of the f e r r i t e , and a d e c r e a s e in r e s i s t a n c e to c r a c k p r o p a g a t i o n due to a c o a r s e n i n g of the alloy c a r b i d e s and a c o a r s e n i n g of the f e r r i t e . In a d d i t i o n it would be e a s i e r to p r o d u c e a p r o p a g a t i n g c r a c k in the c o a r s e r c a r b i d e s .
(c) T h e r m a l s i m u l a t i o n to a peak t e m p e r a t u r e of 1070°C
T h i s s t r u c t u r e , shown in F i g . 6, r e p r e s e n t s that p a r t of the HAZ in which t r a n s f o r m a t i o n f r o m i n h o m o g e n e o u s a u s t e n i t e o c c u r r e d . The i m p a c t p r o p e r t i e s w e r e c o n s i d e r a b l y p o o r e r than in the p r e v i o u s two s p e c i m e n s , due p r i m a r i l y to the p r e s e n c e of u p p e r b a i n i t e and, to a l e s s e r extent, a f u r t h e r c o a r s e n i n g of the f e r r i t e . B o n i s z e w s k i '^^ showed that c o l o n i e s of u p p e r b a i n i t e c o n s i s t i n g of long c r y s t a l s of f e r r i t e s e p a r a t e d by v e r y thin c a r b i d e s , w e r e v e r y s u s c e p t i b l e to cold o r u n d e r b e a d c r a c k i n g . C r a c k i n g o c c u r r e d m a i n l y a c r o s s , but s o m e t i m e s along, the bainitic f e r r i t e g r a i n s .
(d) T h e r m a l s i m u l a t i o n to a peak t e m p e r a t u r e of 1347°C
in which transformation from coarse homogeneous austenite occurred. The impact properties were extremely poor due to the large grain size and the crack susceptibility of upper bainite. The 20 ft. lbs. transition temperature of this structure was raised by about 130 C above that of the parent plate.
4. 4 The Effect of Restraint on the Simulated HAZ Structures
Restraint during thermal simulation had no detectable effect on notch-toughness as Figs. 8 and 9 show. A similar observation was reported for mild steel using the same degree of r e s t r a i n t ^'* and was adequately discussed in that report.
4. 5 The Effect of a Post-Weld Heat Treatment
A subsequent heat treatment to 150°C for 4 hours slightly improved the notch-toughness properties of the simulated weld HAZ structures at low temperatures in all regions investigated. This was probably due to a temp-ering effect on the m i c r o s t r u c t u r e , which would allow some precipitation of carbides from the supersaturated ferrite to take place, making the matrix more ductile and fracture initiation more difficult. The reduction in the notch-toughness values at high temperatures may be due to a coarsening of the carbides reducing the number of c r a c k - a r r e s t o r s .
4. 6 The Importance of the HAZ
By the very nature of the process welding can produce a large number of defects. Such defects, acting at s t r e s s concentrators in an already notch-brittle HAZ. have been the cause of many service failures. Since weld defects cannot be completely avoided, it is important to avoid notch-brittle m i c r o s t r u c t u r e s in the HAZ.
Current philosophy is to aim for m i c r o s t r u c t u r e s with greater resistance to crack initiation since initiation is a necessary step before propagation can occur. The present work showed that, in the welding of QT 35 with a high level of heat input, notch-brittle m i c r o s t r u c t u r e s of upper bainite were produced in the HAZ. The current commercial welding procedure is to use a preheat of 120-150°C and to maintain this preheat until several hours after welding. This treatment not only allows hydrogen to diffuse away from the heated regions, and so reduce the risk of underbead cracking, but also tempers the microstructure by promoting spheroidisation of carbides and precipitation of dissolved carbides. Thus, more ductile microstructures a r e produced.
The C. C. T. diagram for QT 35 (Fig. 16) indicates that if a cooling rate through the transformation temperature range (500°C - 250°C)of not less than 6°C per second is maintained brittle upper bainite will not be formed. Upper bainite is particularly harmful in the region of the HAZ immediately adjacent to the fusion boundary since the structure there is coarse grained and contains the toe of the weld which acts as an appreciable s t r e s s concentrator. In addition high residual s t r e s s e s are often located in this region. The preheat has little effect on the cooling rate through the transformation temperature range but significantly reduces the cooling rate below 250 C thus allowing diffusion of hydrogen to take place.
It is. therefore, suggested that welding procedures for QT 35 should ensure a cooling rate of not l e s s than 6°C per second through the transform-ation temperature range. This could be achieved, in the present instance, by reducing the heat input to the work.
5. Conclusions
1. The weld HAZ of QT 35 was divided into three distinct regions. These were (a) the region of grain coarsening, (b) the region of grain refinement, and (c) the region of partial transformation.
2. A marked increase in hardness occurred in all parts of the visible HAZ. 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. 3. Metallographic examination and hardness measurements showed that the simulated HAZ structures compared directly with those formed in the actual HAZ.
4. The QT 35 parent material simulated to peak temperatures of 788 C, 893°C, 1070°C and 1347°C showed a continuous decrease in notch-toughness with increasing peak temperature of simulation. This was most marked in the material simulated to a peak temperature of 1347°C which had a transition temperature 130°C higher than that of the parent plate. Deterioration of impact properties was directly associated with changes in microstructure. The notch-brittle region produced by simulation to a peak temperature of 1347°C was attributed directly to the presence of a coarse grained upper bainitic microstructure.
5. The effect of an imposed restraint during thermal simulation to produce plastic deformation of 24% had no detectable influence on the final m i c r o -structure and mechanical properties.
6. A post-weld heat treatment at 150°C on structures simulating the partially transformed and grain refined regions produced an improvement in the low temperature notch-toughness properties, but in some cases this was accompanied by a reduction in notch-toughness at high temperatures.
7. A cooling rate through the transformation temperature range of not less than 6^°C per second is necessary to prevent the formation of brittle upper bainite.
6. B i b l i o g r a p h y 1. B a k e r . R. G. et a l . . P r o c . Second C o m m o n w e a l t h Welding C o n f e r e n c e . 1965. 125. 2. B o n i s z e w s k i . T. et a l . . ibid. 117. 3. W a t k i n s o n . F . et a l . Welding in Shipbuilding S y m p o s i u m . I n s t i t u t e of Welding, 1962. 3 1 . 4. N i p p e s . E . F . et a l . Welding J . . 36, 1957, 531 s.
5. N i p p e s , E . F . and Sibley. C. R. . Welding J . . 35. (10). 1956, 473 s. 6. A r n o l d , P . C . . W e l d i n g J . . 3 6 ( 8 ) . 1957. 373 s.
7. S a v a g e . W. F . . and O w c z a r s k i . W. A. , Welding J . . 45. (2). 1966. 55 s. 8. N i s b e t t . H. M. and K e l l y . P . M . , N. C. R. E. R e p o r t No. N. 152, May 1962. 9. O l d r i d g e , D. , D. A. E. T h e s i s , College of A e r o n a u t i c s , Cranfield, 1964. 10. Winn, W. H. , and W o r t l e y , L. , Symp. of Welding in Shipbuilding, Inst.
of Welding. 1962. 36.
11. B e n t l e y , K. P . and Smith. A. A. . P r o c . Second C o m m o n w e a l t h Welding C o n f e r e n c e . 1965. 27.
12. G e o r g e . M. J. and Clifton. T. E. , C o l l e g e of A e r o n a u t i c s Note. C r a n f i e l d , to be p u b l i s h e d .
13. C o w a r d , M. D. and A p p s , R. L. . College of A e r o n a u t i c s Note, M a t . No. 13, C r a n f i e l d , 1967. 14. S m i t h , E. et a l . C o l l e g e of A e r o n a u t i c s R e p o r t , Mat. No. 2, C r a n f i e l d . 15. A n d r e w s , K. W. , J. I. S. I. , 203, (7), 1965. 721. 16. I r v i n e , K. J. and P i c k e r i n g . F . B. , I . S . I . S p e c i a l R e p o r t No. 93, 1965, 120. 17. M o r r i s , A. K. , N. C. R. E R e p o r t No. N156, 1962. 18. Watkinson, F . et a l . P r i v a t e c o m m u n i c a t i o n . 19. B o n i s z e w s k i , T. , B r i t . Welding J . , jl2. (1), 1965, 14. 20. B a k e r , R. G. , P r i v a t e c o m m u n i c a t i o n .
E L E C T R O N MICROGRAPH X 5000 E L E C T R O N M I C R O G R A P H X 10000 HV 225 - 10 F I G . 1 P H O T O M I C R O G R A P H S O F THE S T R U C T U R E O F Q. T . 35 P A R E N T P L A T E
A 700 B X 700
F I G , 2 THE H E A T A F F E C T E D ZONE S T R U C T U R E S ASSOCIATED WITH A B E A D - O N - P L A T E WELD IN Q . T . 3 5 S T E E L
X 700 D
2% N I T A L E T C H 700
in
>
0-5 1-0 l-S 2-0 2-5 3-0
DISTANCE FROM FUSION LINE FIG.3 0.135 HEAT AFFECTED ZONE
HARDNESS SURVEY (LOADSKg.)
UNAFFECTED PARENT PLATE
E L E C T R O N MICROGRAPH 5000 E L E C T R O N M I C R O G R A P H X 10000 + HV 260 - 10 F I G . 4 P H O T O M I C R O G R A P H S O F T H E S T R U C T U R E P R O D U C E D IN Q . T . 3 5 C Y C L E D TO A P E A K T E M P E R A T U R E O F 788°C
E L E C T R O N MICROGRAPH X 5000 E L E C T R O N MICROGRAPH X 20000
+
HV 3 0 0 - 10 F I G . 5 P H O T O M I C R O G R A P H S O F T H E S T R U C T U R E S P R O D U C E D E N Q . T. 35 C Y C L E D TO A P E A K T E M P E R A T U R E O F 893°CE L E C T R O N MICROGRAPH X 6000 E L E C T R O N M I C R O G R A P H X 1 2 0 0 0 H V 3 3 0 t 10 F I G . 6 P H O T O M I C R O G R A P H S O F T H E S T R U C T U R E S P R O D U C E D IN Q. T. 35 C Y C L E D TO A P E A K T E M P E R A T ^ U R E O F 1070°C
— - j y n —
.^)^ i.t*{ E L E C T R O N MICROGRAPH X 5000 éÊÊt E L E C T R O N MICROGRAPH X 5000 HV 370 - 10 F I G . 7 P H O T O M I C R O G R A P H S O F THE S T R U C T U R E S P R O D U C E D IN Q . T . 3 5 C Y C L E D ^ T O A P E A K T E M P E R A T U R E O F 1347°C-20 20 60 100 TESTING TEMPERATURE , °C
FIG. 9 CHARPY V-NOTCH IMPACT (PERCENTAGE CRYSTALLINITY)-TEMPERATURE DATA FOR QT 35 SIMULATED SPECIMENS
60 ^0 20
I
H 80 a: ° ^ Q S cycled us cycicd 60 1.0 20893°C THERMAL CYCLE. 13<7°C THERMAL CYCLE.
0 o as cycled
as cyded
"-«0 -60 -20 20 60 100 UO 180 - 6 0 -20 20 CO 100 ' « 180 TEMPERATURE, ° C
FIG 11. EFFECT OF POST WELD HEAT TREATMENT (C HOURS AT 150° C) ON W B D H A Z PROPERTIES OF a t 35 SIEEL. 788°C THERMAL C V a E . cm—m X'—oo j , o o -B93°C THERMAL CYCLE b -60 1070°C THERMAL C H I E 100 RO 180 TEMPERATLRE,
13i7°C THERMAL CVaE.
o CIS cycled
-100 - 8 0 -60 -«O -20 20 40 60 80 100 C
TESTING TEMPERATURE FIG12 COMPARISON OF CHARPY VNOTCH IMPACT (FTLBS)
-TEMPERATURE CURVES FOR QT35 SIMULATED SPECIMENS
30 50 70 TESTING TEMPERATURE °C
FIG13 COMPARISON OF CHARPY V-NOTCH IMPACT (PERCENTAGE CRYSTALLINITY
tn z < - 4 0 -80 AS REC'D ° DUCTILITY TRANSITION TEMPERATURE AT 20 FT-LBS LEVEL 1600 400 BOO 1200
PEAK TEMPERATURE OF SIMULATION, "c
FIG.14 VARIATION IN TRANSITION TEMPERATURE WITH PEAK TEMPERATURE OF SIMULATION FOR Q.T.35 SPECIMENS
6 0 0 Ui Q : r) < a: 01 0. UJ 500 4 0 0 300 200 ^V i J ' 1 M.. MART ENSITÈ 1 1 1 1 / / / / / /
1 1
\ -.JiTub ' /^ > / / . 0 ,^ j
Ui
(op. r V 111 % TRANSFORMATION s Nc
\ — }wÉR'BAlkl 1 1 1 1 • > - , ^ -TT UPPER BAlt^ltÈ ^ " ' T " * . ^1 "^
^ 1 \m Aus -. 1 ' _ AL-\ d1«r 1 j i J M i l •EN T É ^ _ 1 1 I I 10 100 1000 10000 100000TIME HELD IN CONSTANT TEMPERATURE BATH FROM START OF QUENCH (SECONDS)
FIG15 ISOTHERMAL TRANSFORMATION DIAGRAM FOR Q.T35 AUSTENITISED AT 900'C
10001 800 O o ë 500 => S " 0 0. s. UJ • " 200 0 X ^ \ S s
s
g
S;
\ < ^ ^ ^ ^ ^N
^ HVlS
\ \ L. S >s
N
V ' s % 392 ^s
\ '381_N
^ 372 "H ^n
V s •I 1 ICOOLING RATE FOR THE WELD HEAT AFFECTED ZONE
1 1 1
V
Ê2 320 m
10 10'
TIME TO COOL FROM 1000°C (SECS)
FIG16 CC.T. DIAGRAM FOR Q135 AUSTENITISED AT 1 3 2 5 * 0 "
