CRANFIELD
INSTITUTE OF TECHNOLOGY
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SIMULATED WELD HEAT AFFECTED ZONE
STRUCTURES A N D PROPERTIES OF HY 80 LOW ALLOY STEEL
by
CRANFIELD REPORT MAT. No. 5 February, 1971.
CRANFIELD INSTITUTE OF TECHNOLOGY
SIMULATED WELD HEAT AFFECTED ZONE
STRUCTURES AND PROPERTIES OF HY 80 LOW ALLOY STEEL
by
G. T. B. Kellock, D. A. E. , A . I . M . . C. Eng., A. F . R.Ae.S. E. Smith, P h . D . . B. Sc, , A . I . M .
A. R. Sollars, B. Sc.
SUMMARY
Single and double thermal cycle simulation of heat affected zone (HAZ) structures has been used to study the structural and property changes produced by submerged arc welding of HY 80 steel. The effectiveness of the temper bead technique and post-weld heat treatment at 650°C have also been examined,
Marked degradation of impact properties occurs, particularly in the grain-coarsened region of the HAZ. Misalignment of a temper bead may result in a further impairment of notchtoughness. Postweld heat t r e a t -ment is shown to be effective in restoring properties to levels similar to those of the parent plate.
The significance of the r e s u l t s is discussed in relation to submarine applications.
1. INTRODUCTION 1 2. E X P E R I M E N T A L 2 2. 1 M a t e r i a l s 2 2 . 2 P r o c e d u r e 3 3. R E S U L T S 5 3. 1 P a r e n t m a t e r i a l and HAZ 5 3 . 2 P a r e n t p l a t e banding 6 3. 3 Simulated weld HAZ s t r u c t u r e s 7
4. DISCUSSION 4. 1 Simulated weld t h e r m a l c y c l e s 9 4. 2 I n t e r p r e t a t i o n of C h a r p y d a t a 10 4 . 3 Notch t o u g h n e s s r e q u i r e m e n t for HY 80 in s u b m a r i n e s 11 4 . 4 Significance of the r e s u l t s 12 5. CONCLUSIONS 15 R E F E R E N C E S 17 TABLES 1. M e c h a n i c a l p r o p e r t y d a t a for a s - c y c l e d s i m u l a t e d s p e c i m e n s 20 and the p a r e n t m a t e r i a l 2. M e c h a n i c a l p r o p e r t y d a t a for s i m u l a t e d s p e c i m e n s , post c y c l e 21 h e a t t r e a t e d at 650 C. F I G U R E S
1. T h e r m a l c y c l e s p r o d u c e d in the p a r e n t plate adjacent to the weld
2. P a r e n t p l a t e m i c r o s t r u c t u r e s
3. Heat affected zone h a r d n e s s s u r v e y
4. The n a t u r e of banding in the p a r e n t p l a t e , a f t e r t h e r m a l c y c l i n g to a peak t e m p e r a t u r e of 930°C.
5. Effect of t h e r m a l cycling and s u b s e q u e n t p o s t - c y c l e h e a t t r e a t m e n t on C h a r p y V - n o t c h i m p a c t data
6. Effect of s u b s e q u e n t t h e r m a l c y c l i n g and p o s t - c y c l e h e a t t r e a t m e n t on C h a r p y V - n o t c h i m p a c t d a t a for s p e c i m e n s i n i t i a l l y c y c l e d to a peak t e m p e r a t u r e of 1275°C.
7. Effect of s u b s e q u e n t t h e r m a l c y c l i n g and p o s t - c y c l e h e a t t r e a t m e n t on C h a r p y V notch i m p a c t d a t a for s p e c i m e n s i n i t i a l l y c y c l e d to a peak t e m p e r a t u r e of 930°C.
8. Effect of subsequent thermal cycling and post-cycle heat treatment on Charpy V-notch impact data for specimens initially cycled to a peak temperature of 765°C.
9. Charpy V-notch impact (% crystallinity) data for specimens initially cycled to 12750C.
10. Charpy V-notch impact (% crystallinity) data for specimens initially cycled to 930°C and 765°C.
11. Simulated weld HAZ structures produced by single and double cycles, with initial cycling to 1275°C.
12. Simulated weld HAZ structures produced by single and double cycles, with initial cycling to 930°C.
13. Simulated weld HAZ structures produced by single and double cycles, with initial cycling to 765°C.
14. Post-cycle heat treated structures of single and double cycled specimens with initial cycling to 1275°C.
15. Post-cycle heat treated structures of single and double cycled specimens with initial cycling to 930°C.
After the end of World War II, much effort has been ex|)ended in the development of high notch-toughness steels, particularly for naval construction. Considerable experience had been obtained in the marine industry in using the quenched and tempered STS steel. Test results on low carbon content heats of this material prompted modification to a low carbon, nickel, chromium, molybdenum steel which, when quenched and tempered, developed a yield strength of 550 N/mm.^ (80,000 Ibf/in^) in thicknesses up to 32 mm, with notch-toughness values, based on the Charpy V-notch impact test, of 68J (50 ft. Ibf) at -84°C. With composition variation to ensure good hard-enability in different plate thicknesses this steel was referred to a s HY 80 and produced in plate form to U. S. military specifications MIL - S - 16216D, E and F , (Refs. 13). The present U.S. military specification, MIL S -16216G (Ref. 4), quotes only one composition range for all plate thicknesses.
In the application of nuclear power to submarine propulsion the hull diameter had to be increased over that required for conventional power, in order to encase the necessary equipment. Due to its extremely good strength : weight ratio, notch toughness at low temperatures and anti-ballistic properties when compared with other steels, HY 80 has been extensively used in
fabrication of p r e s s u r e hulls and other components for the U. S. nuclear submarine fleet. In the U.K., it is now superseding the British QT 35 quenched and tempered low alloy steel in the production of the Dreadnought class of nuclear powered submarines. The present Ministry of Defence specification DG/SHIPS/PS/9027, (Ref. 5) requires plates to the U.S. military specification MIL - S - 16216G, (Ref. 4). McKee, (Ref. 6) has discussed the position of HY 80 steel in relation to submarine design practice and Heller et al,
(Ref. 7) in an excellent review paper, have evaluated its use as a structural material for nuclear powered submarines.
Since welding is extensively used in submarine fabrication it is important to be aware of its effect on the structure and properties of the parent material. In low alloy steels the fracture toughness of the heat affected zone (HAZ) may be severely impaired by welding, resulting in increased susceptibility to brittle fracture and possible cold cracking.
(Refs. 8-10). F r a c t u r e toughness of the weld HAZ i s dependent upon the types of structures produced by the weld thermal cycles and this is controlled to some extent in HY 80 steel by limiting the heat input to the range 1.2 -2. -2. k J / m m (30 - 55kJ/in. ) and using a preheat and interpass temperature in the range 120°C-150°C, (Ref. 5). These conditions favour the formation of structures consisting predominantly of autotempered martensite with perhaps some lower and upper bainite. In this class of steels this type of structure has been shown to possess better fracture toughness properties than those produced by heat inputs above this range, which contain a greater proportion of upper bainite, (Ref. 11). However, even when the heat input and preheat temperature requirements are adhered to, there is still some doubt as to whether the notch toughness of the weld HAZ provides adequate resistance to brittle fracture during service. In view of the fact that HY 80 has been used in marine construction since 1952, it is rather surprising that very little infornaation regarding its structure as produced by standard heat treatment procedures, let alone the complexities introduced by the weld thermal cycle, has been published. Apart from the work of Dolby (Ref. 12) on weld HAZ structures produced by heat inputs lying outside the specified range, no electron microscope study of HY 80 steel has been reported
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-in the literature.
Current U. 5. Navy and Ministry of Defence (Navy) specifications call for the application of the tempering bead technique for all multi-run welds in HY 80 steel. Its object is to reduce the high hardness and relatively low fracture toughness associated with undesirable structures in the HAZ,
particularly in the region of the toe of the weld. The technique is of primary importance in respect of the final surface runs since sub-surface layers generally have their HAZ tempered by subsequent layers. The positioning of the tempering bead must be carefully controlled, since if it is not centered properly with respect to the two HAZ boundaries in contact with the surface, a significant tempering effect cannot be obtained in these parts of the HAZ. In fact, it is possible to envisage a multiplicity of combinations of thermal cycles in the edge bead HAZ due to the application of a tempering bead. In an effort to minimise the effect, the welding specification for HY 80 steel requires that approximately 3. 2mm of the edge beads should be exposed after deposition of the temper bead. Oldridge (Ref. 13) showed that there are such immense practical difficulties involved in maintaining accurate positioning of the tempering bead during deposition that its efficiency in providing a tempering effect in the HAZ of the last deposited layers, adjacent to the parent material, may be considered to be open to conjecture.
The Ministry of Defence specification (Ref. 5) considers the post-weld heat treatment of HY 80 weldments for the purpose of s t r e s s relief to be neither necessary nor desirable. If, for any reason, s t r e s s relief is required it is undertaken at 550°C jl 15°C, this temperature being held for 1 hour per inch of thickness of the thickest member of the weldment. The potential benefit to be derived from post-weld heat treatment in t e r m s of restoration of good notch toughness in the HAZ appears to have been
neglected in the past. It is the authors' opinion however that such treatment could be very beneficial to the weld HAZ and perhaps also to the weld metal, and thus remove the uncertainty concerning the effectiveness of the tempering bead technique.
This present r e s e a r c h has a three-fold object
:-(a) to establish for a given heat input condition and preheat temperature the nature of the weld HAZ due to a single-run weld.
(b) by employing double cycle simulation techniques using the same welding conditions a s in (a), to indicate the general nature of tempering in a pre-existing HAZ due to the heat flow from a subsequent weld run in a multi-run weldment and to predict the effectiveness of the temper-bead technique.
(c) to ascertain the effectiveness of a post-weld heat treatment at 650^0 for 1 hour on the structures produced by (a) and (b).
2. EXPERIMENTAL 2. 1. . Materials
which c o n f o r m e d to specification MIL - S-16216G
c
0 . 1 6 Mn 0 . 3 2 Si 0.30 S 0 . 0 1 7 P 0, 006 Ni 2 . 5 4 C r 1.31 Mo 0. 28 V 0 . 0 1 Al 0. 015 N 0 . 0 1 T e n s i l e and C h a r p y V - n o t c h i m p a c t t e s t s conducted on s p e c i m e n s obtained f r o m the plate m i d - t h i c k n e s s yielded the following r e s u l t s , which c o n f o r m e d to specification MIL - S - 16216G. Specimen O r i e n t a t i o n T r a n s v e r s e Longitudinal C h a r p y E n e r g y * at - 8 4 ° C J 107 127 ft. Ibf 79 93 Yield Strength N / m m 576 594 tonf/in 3 7 . 4 38. 6 U . T . S . N / m m 719 735 t o n f / i n 4 6 , 7 4 7 . 8 * notched in the t h r o u g h t h i c k n e s s d i r e c t i o n 2. 2. P r o c e d u r e 2, 2. 1. P r e p a r a t i o n and e x a m i n a t i o n of weld.A b e a d - o n - p l a t e weld w a s p r o d u c e d by a subnaerged a r c welding unit using 4. 7 6 m m (0. 1875 in) d i a m e t e r m i l d s t e e l f i l l e r w i r e with the following c o n d i t i o n s : -C u r r e n t (amps) 500 Voltage 30 T r a v e l Speed m m / s e c 7 . 0 5 P r e h e a t T e m p . (°C) 120 Heat Input k J / m m 2 . 1 3
A t r a n s v e r s e s e c t i o n t h r o u g h the weld HAZ w a s e x a m i n e d u s i n g o p t i c a l m e t a l l o g r a p h y , and h a r d n e s s d e t e r m i n a t i o n s w e r e m a d e at 0. 5mm i n t e r v a l s ,
s t a r t i n g f r o m the fusion b o u n d a r y , using a Zwick h a r d n e s s t e s t e r and l o a d s of 0 . 5 kg and 5 kg.
2. 2, 2. Simulation of weld HAZ s t r u c t u r e s
Weld t h e r m a l c y c l e s i m u l a t i o n w a s c a r r i e d out on an a p p a r a t u s
d e s i g n e d and built at Cranfield in which a m a t e r i a l blank of s u i t a b l e d i m e n s i o n s i s h e a t e d by v i r t u e of i t s own r e s i s t a n c e to the p a s s a g e of an e l e c t r i c c u r r e n t and cooled by the flow of w a t e r t h r o u g h hollow b r a s s c l a m p i n g b l o c k s . The C r a n f i e l d s i m u l a t o r h a s b e e n d e s c r i b e d in d e t a i l by Clifton and G e o r g e , (Ref, 14).
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-10. 7mm X -10. 7mm x -10. 7mm x 83mm and were machined from the mid-thickness of the plate axially transverse to the rolling direction. Coward, (Ref. 15), showed that the application of restraint by rigid clamping of the blanks during simulation of a similar steel did not affect the performance in the Charpy V-notch impact test. In this present study all blanks were allowed axial freedom of movement during thermal cycling.
Four thermal cycles with peak temperatures of 1275°C, 930°C,
765°C and 650°C, corresponding to those experienced by the grain coarsened, grain refined, intercritical, and subcritical regions of the weld HAZ
respectively, were used for simulation. The 1275°C and 7650C peak temperature cycles were measured directly by Smith (Ref. 16) in the HAZ of a submerged arc bead-on-plate weld using the same set of conditions described in section 2. 2. 1. The technique has been described by Coward, (Ref. 15). The 930°C and 650°C peak temperature cycles were computed by Kellock (Ref. 17) from a s e r i e s of programs incorporating the results of Smith, (Ref. 16). These programs included functions to account for the release of latent heat from the solidifying weld pool and the variation of thermal conductivity with temperature. The computed thermal cycles correlated well with the measured cycles. The four thermal cycles are shown in Fig. 1.
The following thermal cycles and combinations of thermal cycles were used to examine the HAZ structures of single and multipass
weldments:-Peak Temperature, °C F i r s t Cycle 1275 930 765 1275 1275 Second Cycle 1275 930 Peak Temperature, °C 1 F i r s t Cycle 1275 1275 930 930 765 Second Cycle 765 650 765 650 650
2. 2, 3. Post-cycle heat treatment
Half the total number of simulated blanks were given a post-cycle heat treatment at 650°C for 1 hour, the specimens being placed in the furnace at 250°C and heated to 650°C in 50 minutes. After one hour at temperature the specimens were removed from the furnace and allowed to cool in still air.
2. 2. 4, Mechanical testing of simulated HAZ structures
Ten standard Charpy V-notch impact specimens were prepared from the simulated blanks for each condition studied with the notches machined in
the control thermocouple position and in the through thickness direction of the original plate. Impact transition curves were determined by testing between - 196°C and 40°C. A mixture of iso - pentane and liquid nitrogen was used for t e s t s below room temperature and hot water for t e s t s above room temper-ature. Test temperatures were estimated to be accurate to 't 3°C.
Three Hounsfield tensile test specimens with a gauge length of 7. 6mm and 4. 5mm diameter were prepared from the simulated blanks for each con-dition, with the gauge length accurately positioned within the thermally cycled zone at the centre of each blank. The specimens were tested on a standard Instron tensile machine using a strain rate of approximately 3 x 10 " sec. " . The values of 0. 2% proof s t r e s s , U. T. S. , and reduction of area were recorded.
One simulated blank from each condition was sectioned through the centre of the thermally cycled zone and the hardness determined using the Zwick hardness t e s t e r with a load of 5kg.
2. 2. 5. Metallographic examination of simulated HAZ structures
The sections used for hardness determinations were subsequently prepared for optical metallographic examination, all structures being
successfully etched in 2% nital. P r i o r to the preparation of carbon extraction replicas for electron microscopy, the specimens were given a further etch in 2% nital in order to ensure the production of satisfactory replicas. The replicas were extracted with a 10% nital solution and examined in a Siemens Elmiskop lA electron microscope.
2. 2. 6. Examination of banding
Band width measurement and 0. 5kg load hardness surveys were made on specimens which had been simulated using the 930°C peak temperature cycle. Specimens in this condition were chosen because they gave the best inter-band optical contrast after etching. Electron probe microanalysis was c a r r i e d out at the Welding Institute to determine the nature of the m i c r o s e g r e -gation within the HY 80 steel plate.
3, RESULTS
3, 1. Parent material and weld HAZ
The parent plate m i c r o s t r u c t u r e . Fig. 2, was typical of a quenched and tempered low alloy steel in that it contained a dispersion of carbides in a ferrite matrix. Three carbide types are apparent in Fig. 2. There is a Widmanstatten a r r a y of fine rod-shaped particles situated within ferrite sub-grains, having a length of approximately 0. lium. In addition spheroidised carbides approximately 0. 5^m in diameter and some larger, more elongated particles about l/.im long are situated mainly at ferrite sub-boundaries. Dolby (Ref. 18) has indicated that the grain boundary carbides consist of both M7C3 and Fe3C while the fine intragranular precipitates are invariably Fe C.
6
-distinct regions
:-(a) the grain coarsened region extending up to approximately 0. 6mm from the fusion boundary,
(b) the grain refined region, between 0. 6mm and 2. 5mm from the fusion boundary, and
(c) the intercritical region, between 2.5mm and 3.0mm from the fusion boundary.
Due to the high temperatures experienced within the grain coarsened region, rapid austenite grain growth occured during thermal cycling but the austenite grain size decreased, with decreasing peak temperatures, as the distance from the fusion boundary increased. Within the prior austenite grains the structure was acicular but the details could not be resolved using optical microscopy. The boundary separating the grain coarsened and grain refined regions is rather diffuse. In the grain refined region the structure consisted of a fine irresolvable ferrite-carbide aggregate with a fine, and decreasing, prior austenite grain size. The intercritical region contained an increasing proportion of untransformed ferrite as the distance from the fusion boundary increased and the transformed regions appeared to consist of a fine carbide aggregate, once more not capable of being resolved using conventional light microscopy.
The r e s u l t s of the 5kg, and 0. 5kg. load hardness t r a v e r s e s are shown in Fig. 3. The 5kg. r e s u l t s show a continuing increase in hardness on moving towards the fusion boundary reaching a maximum value of
430 HV5 at the fusion boundary. Large fluctuations superimposed on the same general trend were obtained with the 0. 5kg. load t e s t s , the fluctuations being particularly noticeable in the grain refined region. This is thought to be due to the banded structure of the parent plate resulting from alloying element segregation.
3. 2. Parent plate banding
The existence of banding was not very evident in the parent plate or after cycling to a peak temperature of 1275°C. However, after cycling to 765°C and 930°C, banding was very noticeable at low magnifications. Banding in a 930°C peak temperature cycled specimen a s revealed by etching in 2% nital is shown in Fig. 4. Measurement of band width and hardness gave the following r e s u l t s :
-Band type Light etching Dark etching Band width (mm) Mean 0. 056 0. 195 Std. Dev. 0. 031 0.044 HV 0. 5 Mean 436 311 Std. Dev. 7 11 HV5 1 Mean 353
Std. Dev.
1 17E l e c t r o n probe m i c r o a n a l y s i s of t h r e e t y p i c a l bands gave the following r e s u l t s : -E l e m e n t M a n g a n e s e N i c k e l Chromium: Molybdenum Weight % in band D a r k etching 0. 41 2. 20 1. 70 0. 20 Light etching 0. 59 3. 50 2 . 4 0 0. 26 D a r k etching 0 . 3 9 2 . 3 1 1.20 0. 18
3. 3. Simulated weld HAZ s t r u c t u r e s
The C h a r p y V-notch i m p a c t t e s t r e s u l t s on single and double c y c l e d s i m u l a t e d s p e c i m e n s , with and without p o s t - c y c l e heat t r e a t m e n t , a r e given in F i g u r e s 5 - 1 0 . F i g u r e s 5 - 8 show the e n e r g y a b s o r p t i o n - t e m p e r a t u r e c u r v e s and F i g u r e s 9 and 10 the f r a c t u r e a p p e a r a n c e - t e m p e r a t u r e c u r v e s . F i g u r e 5 a l s o i n c l u d e s the e n e r g y t r a n s i t i o n c u r v e s for the p a r e n t plate in both the t r a n s v e r s e and longitudinal d i r e c t i o n s .
The r e s u l t s obtained f r o m t e n s i l e and h a r d n e s s t e s t i n g a r e given in T a b l e s 1 and 2 for the a s - c y c l e d and p o s t - c y c l e d heat t r e a t e d c o n d i t i o n s r e s p e c t i v e l y . Data for the p a r e n t p l a t e i s a l s o included in T a b l e 1. The s t r u c t u r e s p r o d u c e d by the s i m u l a t e d weld t h e r m a l c y c l e s and p o s t - w e l d heat t r e a t m e n t a r e shown in F i g u r e s 11 - 15.
3. 3, 1. Single c y c l e s i m u l a t i o n
A single cycle to a peak t e m p e r a t u r e of 1275°C p r o d u c e d a s t r u c t u r e c o n s i s t i n g p r e d o m i n a n t l y of a u t o - t e m p e r e d lath m a r t e n s i t e . F i g . 11a, with
s m a l l a m o u n t s of u p p e r and l o w e r b a i n i t e . F i g . l i b . T h e r e w a s a c o n s i d e r a b l e i n c r e a s e in proof s t r e s s , U . T . S . and h a r d n e s s and d e c r e a s e in ductility when c o m p a r e d with p a r e n t plate (Table 1), while the C h a r p y v a l u e s w e r e s e v e r e l y affected (Fig. 6).
After a single c y c l e to a peak t e m p e r a t u r e of 930 C s l i g h t l y b e t t e r e n e r g y a b s o r p t i o n w a s obtained. F i g u r e 12a shows a t y p i c a l s t r u c t u r e c o n s i s t i n g of a v e r y fine f e r r i t e - g l o b u l a r c a r b i d e a g g r e g a t e and a r e a s of a u t o - t e m p e r e d m a r t e n s i t e . In F i g . 12b the halo affect a r o u n d s o m e c a r b i d e p a r t i c l e s s u g g e s t s that not a l l the c a r b i d e w a s t a k e n into solution during the t h e r m a l c y c l e . The m e c h a n i c a l p r o p e r t i e s , T a b l e 1 , show a c o n s i d e r a b l e i n c r e a s e in proof s t r e s s , U . T . S . and h a r d n e s s and d r o p in d u c t i l i t y c o m p a r e d with p a r e n t p l a t e . The C h a r p y p r o p e r t i e s . F i g s . 5 and 10a show a c o n s i d e r -able d r o p in e n e r g y a b s o r p t i o n c o m p a r e d with p a r e n t p l a t e .
A single cycle of 765 C r e s u l t e d in p a r t i a l a u s t e n i t i s a t i o n . A high d e n s i t y of c o a r s e c a r b i d e w a s evident in sonae a r e a s of the s t r u c t u r e . F i g . 13a,
Many of the fine rod-shaped carbides present in the parent plate were taken into solution and retained there on cooling. The nature of the transformed a r e a s was difficult to identify but they could have been bainitic or martensitic formed during the cooling of carbon-enriched austenite. There was a small increase in proof s t r e s s , U . T . S . and hardness (Table 1) although the Charpy properties were similar to those of the parent plate in the transverse
direction (Fig. 8).
3. 3. 2. Double cycle simulation
A second cycle to 1275°C after an initial 1275°C cycle produced only slight changes in structure. The prior austenite grain size was somewhat larger and more well-defined upper bainite colonies were present. The m.ost significant change in mechanical properties was the resultant fall in upper shelf energy from 50J to 37J(37 ft. Ibf. to 27 ft. Ibf. ) as shown in Fig. 6. A second cycle to 930°C peak temperature after an initial 1275°C cycle produced a structure composed mainly of auto-tempered martensite (Fig. l i e ) in which the laths were much finer than the structures produced by single or double cycling to 1275°C peak temperature. Some upper bainite was also present. The prior austenite grain size was of duplex type, showing a larger grain size due to the initial 1275°C cycle within which was a much finer one due to the lower temperature second cycle. The proof s t r e s s , U . T . S . and hardness were increased slightly by the second cycle (Fig. 6) probably due to grain refinement.
A noticeable improvement occurred when a second cycle to a peak tenaperature of 765°C followed the initial 12750C cycle. Fig. l i d shows a virtually continuous constituent in the prior austenite grain boundaries due to the initial cycle. As seen in Fig. l i e , this constituent is composed of small autotempered martensite units. The a r e a s surrounded by this matrix con-sisted of ferrite and course carbide. There was a marked drop in proof s t r e s s , U, T, S, and hardness (Table 1) and some improvenaent in Charpy performance (Figs, 6 and 9b).
Following an initial 1275°C cycle with a cycle to 650°C peak produced a tempering effect. Carbides precipitated at autotempered
martensite lath boundaries (Fig. 14). Original carbides in the autotempered martensite show some degree of growth. Precipitation had also occurred within the martensite laths and in the prior austenite grain boundaries. The proof s t r e s s , U . T . S . and hardness were lower than after the initial 1275°C cycle and Charpy performance improved, (Pigs. 6 and 9b). It should be noted, however, that this second cycle produced slightly higher strengths than a second cycle to 765°C.
When an initial cycle to 930°C peak was followed by a cycle to 765°C, only partial austenitisation occurred and the resulting structure (Fig. 12c) consisted of fine-grained ferrite and regions of a fine ferrite-carbide dispersion. There was also some evidence of only partial solution of carbides. The mechanical properties showed a marked drop in proof s t r e s s , U . T . S . and hardness (Table 1) and a slight improvement in Charpy performance (Figs. 7 and 10a).
A second cycle to 650 C after a 930°C cycle produced a tempering effect as shown by an increased density of carbides, although the distribution was inhomogeneous (Fig. 12d). The strength properties were higher than for the parent plate in the transverse direction while the energy absorption
properties, although improved by the second thermal cycle, were still markedly inferior to the parent plate (Fig. 7).
Specimens initially cycled to 7650C and subsequently cycled to 650°C showed structures resembling the parent plate (Fig. 13b) probably due to r e -precipitation of carbides taken into solution during the initial cycle. The presence of small rod-shaped particles in some regions (Fig. 13c) indicates that some auto-tempered martensite was produced by the first cycle. The mechanical properties were very similar to the transverse properties of the parent plate.
3. 3. 3. Post-cycle heat treatment
The post-cycle tempering at 650°C for 1 hour produced structures consisting essentially of temper carbides in a ferrite matrix. As for the parent plate, three types of carbides were identifiable viz. fine rod shaped particles and c o a r s e r globular and elongated carbides. The distribution of these types varied within groups of specimens.
(a) Specimens initially cycled to 1275°C generally contained fine rod-shaped and coarse elongated carbide particles with only a small
proportion of globular particles. The coarse elongated carbides were sited mainly at prior austenite grain boundaries and between individual ferrite laths, a s shown in Fig. 14a for a specimen originally given a single cycle to 1275°C, Fine rod shaped particles occurred in the tempered martensite laths as shown in Fig. 14b for a specimen given two cycles to 1275°C peak. The structure produced by tempering after a second cycle to 765°C is
particularly interesting. F i g s . 14c and 14d, Martensitic a r e a s delineating original prior austenite grain boundaries have been tempered to produce a network of dense, mainly elongated carbides, surrounding a r e a s containing Utile precipitate. The proof s t r e s s , U . T . S . and hardness of all these structures were higher than the parent plate (Table 2), and a considerable improvement in Charpy performance resulted, (Figs, 6 and 9b).
(b) Specimens initially cycled to 930°C contained mainly globular carbides, with some elongated carbide particles (Figs. 16a and 16b). The temper carbides were fairly uniformly distributed throughout a ferrite matrix and tensile, hardness and Charpy properties were close to those of the parent plate (Table 2, Figs. 7 and 10a).
(c) Specimens initially cycled to 765°C and tempered at 650°C showed structures very similar to those of the parent plate and the mechanical properties were also very similar (Figs. 8 and 10b),
4. DISCUSSION
4. 1. Simulated weld thermal cycles
10
-with those produced by the simulation equipment was generally very good. However, difficulty was experienced in reproducing the sharp peak in the 1275°C cycle, the electrical characteristics of the simulation equipment producing a broader peak. This could result in the simulated samples showing a slight increase in the prior austenite grain size and a more homogeneous austenite compared with the corresponding region in a weld heat affected zone.
Inflections in the cooling curves were observed for most of the simulated specimens in the temperature range 550°C - 400°C (Ref. 17) and have been attributed by Inagaki et al (Ref. 19) to the exothermal transform-ation of austenite.
4. 2, Interpretation of Charpy data
The successful performance of a welded steel structure in avoiding brittle fracture will depend upon the fracture toughness
characteristics of the parent material, the HAZ and the weld metal. Low fracture toughness in any of these regions may lead to complete failure of the structure by catastrophic brittle fracture since each of these regions provide a continuous path along which such fractures may propagate.
The Charpy test has been the most widely used method of assessing fracture toughness in structural steels. Interpretation of the results can, however, be misleading unless it is based on a vast amount of experience or else correlated with some other test procedure which gives a more realistic appraisal of service performance. This is due partly to the fact that the energy measured by the test does not permit conditions for initiation and propagation to be distinguished or the critical s t r e s s and strain p a r a m e t e r s responsible for fracture to be separated. The test also uses impact loading conditions and thus the results will not be directly applicable to static loading conditions.
Pellini et al (Refs. 20 - 24), however, have developed a procedure for the fracture-safe design of steel structures based on a fracture analysis diagram which takes account of flaw size, s t r e s s level and service temperature. The procedure defines three critical transition t e m p e r a t u r e s :
(a) The nil-ductility transition (NDT) temperature below which the steel loses all ability to deform in the presence of a sharp crack.
(b) The fracture transition elastic (FTE) temperature below which brittle fracture will run through elastically loaded material.
(c) The fracture transition plastic (FTP) temperature above which only shear tearing is possible irrespective of severity of plastic loading.
The temperature intervals between these critical transition temperatures a r e remarkably similar for a wide range of steels including mild steels, high strength steels, cast steels, and 12% Cr steels (Ref. 21). so that it is possible to establish the NDT temperature by the use of the explosion bulge test. The proposed relationships between the three critical
transition temperatures are as follows:
FTE = NDT + 33°C (+ 5°C) F T P = NDT + 72°C (t 10°C)
Pellini et at (Refs. 20 - 24) have shown that a correlation can be established between the NDT temperature measured by the drop weight or explosion bulge tests and the Charpy energy level at this temperature. It must be emphasised, however, that this correlation is unique and is only valid for other steels of similar composition, manufacture, and heat t r e a t -ment. When this correlation has been established Charpy data can be used to predict the three critical transition temperatures and this enables limit-ations to be placed on the steel which, if adhered to, will ensure that brittle fracture does not occur.
For welded structures the concept of inherent flaws is considered to be quite realistic since they may arise from lack of penetration, lack of fusion, slag inclusion, porosity or cracking. Pellini and Puzak (Ref. 23) have provided extensive failure and structural test data to support the validity of the procedure.
4. 3. Notch toughness requirement for HY80 in submarines
Pellini and Srawley (Ref. 22) have discussed the notch-toughness requirements of HY80 in submarine structures, two types of structure being considered. For non-combatant submarines the lowest service temperature should be above the FTE temperature, so that fracture a r r e s t protection is provided even for parts of the structure where tensile s t r e s s e s are close to yield point levels. For combatant submarines which are required to with-stand explosive loading, resulting in extensive plastic deformation of the hull structure, the F T P temperature is required to be below the lowest service temperature.
For HY80 parent material it was established that the FTE and F T P temperatures are approximately -68°C and-40°C respectively and that the NDT temperature (-90^0) c o r r e l a t e s well with the 40-45 ft. lb Charpy temperature (Ref. 21). Winn (Ref. 25) reported a similar correlation for QT 35 on specimens cut in the rolling direction and this has been confirmed in recent work at Cranfield (Ref. 26). Thus, the fracture toughness
requirements for the two types of submarines described above can be stated a s follows, based on safe operation at -10 C which is below the lowest temperature encountered during submerged operation
:-(a) 68J (50 ft. lb. ) minimum at -40°C for non-combatant submarines (b) 68J (50 ft. lb. ) minimum at -68 C for combatant submarines.
A similar attempt to evaluate the behaviour of the HAZ is much more difficult. It has been suggested that fractures which initiate in the HAZ will propagate into either parent material or weld metal since the welds are normally of V or double V geometry and thus produce a HAZ slanted 45 - 60 C with respect to the usual s t r e s s vector.
- L
Therefore, protection against fracture initiation only is indicated. However there is contradictory evidence (Ref. 33) that propagation can occur solely through the HAZ in such a preparation. Nevertheless, because s t r e s s -relieving is not normal practice for submarine structures, high yield residual s t r e s s e s will be present in the HAZ so that operation above the NDT temperature is required to protect against fractures initiating from small flaws for a wide range of nominal s t r e s s e s .
At the present moment it is not possible to state categorically what the notch toughness requirements are in t e r m s of Charpy data since the correlations for the types of structure produced in the HAZ have not been determined. However, Puzak andPellini (Ref. 21) have shown that for structures of relatively high hardness which have a maximum Charpy energy of about 41J (30 ft. lb) or l e s s , the NDT temperature is generally related to the 14 - 20J (10 - 15 ft. lb. ) Charpy temperature. Recent work at Cranfield (Ref. 26) on plates of QT35 steel quenched to produce structures similar to those found in the grain coarsened region of the HAZ indicates that the NDT temperature is equivalent to the 23-31J (17-23 ft. lb.) Charpy temperature. Thus, it would appear that the fracture toughness requirement for grain coarsened HAZ structures in HY80 will be in the region of 34J (25 ft. lb. ) at -10°C.
4.4. Significance of the r e s u l t s 4. 4. 1, Parent material
The HY80 used in the present work easily met the 68J (50 ft. lb. ) minimum Charpy energy at 68°C considered necessary for the more severe conditions required of combant submarines. Fig. 5 shows that at this temperature the energy levels were approximately 149 and 115J (110 and 85 ft. lb. ) respectively for the longitudinal and transverse directions. It is of interest to conapare these values with those reported for QT 35 since HY80 has recently replaced this steel in the U.K. submarine programme. Smith and Apps (Refs. 26 and 27) have determined equivalent Charpy energy values at -60°C of 90 and 61J (66 and 45 ft. lb, ) respectively for the
longitudinal and transverse directions of QT35. Thus, the replacement of QT35 by HY80 in the U.K. would appear to be justified. Dolby (Ref. 12) has reported a crack opening displacement transition temperature some 20°C lower for HY80 compared with QT35. The superior notch-toughness of HY80 has been attributed to the following factors (Ref. 18)
:-(a) a smaller grain size giving greater resistance to crack initiation and propagation
(b) the higher Ni content of HY80 which improves the basic properties of the ferrite
(c) an increased density of carbides and different size distribution increasing resistance to crack propagation
(d) A lower dislocation density within the subgrains
in that the light etching bands were found to contain a higher concentration of aUoying elements than the dark etching bands. This is in direct contrast to another examination of banding in HY80 (Ref. 18), in which the dark etching bands were found to be richer in alloy content. However, in the present work banding was examined in material that had been thermally cycled above the upper critical temperature so that the apparent anomaly can be explained by the fact that the alloy-enriched bands transformed to martensite during Cooling and thus appeared lighter than the alloydenuded bands which t r a n s -formed to a higher temperature transformation product containing more carbide precipitates.
4 . 4 . 2 . Single pass weld HAZ
The single cycle simulation technique enables a study to be made of the HAZ in single pass welds which also may be considered to represent the worst case for multi-pass welding where HAZ hardnesses approaching those in single pass welds can be found in the final edge bead HAZ.
The results. Fig. 6, show that the transformed HAZ has significantly lower energy absorption and upper shelf energy compared with the parent plate, these changes becoming more marked on approaching the fusion boundary. A similar pattern of results has been reported on other quenched and tempered low allow steels (Refs. 11, 15, 27-30), Dolby (Ref, 12) has also shown a significantly reduced resistance to fracture initiation in the transformed HAZ of HY80. This reduction in fracture toughness is accompanied by marked increases in proof s t r e s s and hardness (Table 4). The high hardness associated with predominantly martensitic structures, particularly in the grain coarsened HAZ, indicates that precautions need to be taken to keep hydrogen to a minimum in order to avoid HAZ cold cracking.
The Charpy data for the grain coarsened HAZ shows an energy level slightly in excess of 41J (30 ft. lb. ) at -10 C which is above the 34J (25 ft, lb. ) level considered necessary to indicate an NDT temperature below -10 C, Thus, the probability of fractures initiating from small flaws in the HAZ would appear to be small although more work is necessary to confirm the correlation between the NDT temperature and the Charpy energy level at this temperature and to investigate the variation in properties
between different casts of HY80.
A similar investigation into the weld HAZ properties of QT35 steel (Ref. 27) indicated a Charpy energy level of about 30J (22 ft. lb. ) at -10 C for the same welding conditions used in the present work and would thus suggest that QT35 is border-line with respect to NDT requirement. This would appear to be further confirmation of the supjeriority of HY80 over QT35. Dolby (Ref. 12) also showed a lower resistance to fracture initiation in the grain coarsened HAZ for QT35 compared with HY80 when initiation occurred by quasi-cleavage although the position was reversed at higher temperatures when the initiation mode was by stable tearing. The
superiority of HY80 over QT35 can be attributed to the increased hardenability of HY80 giving predominantly martensitic structures in the t r a n s -formed HAZ, the increased Ni content improving the basic properties of the structure, and the smaller prior austenite grain size caused by the pinning action of aluminium nitride particles. However, there is evidence
14
-to show that the lower Mn/S ratio and the higher Ni content of HY80 increases the susceptibility to hot cracking, (Ref. 18).
The improved Charpy performance observed in the grain refined HAZ compared with the grain coarsened HAZ in HY80 can be attributed to the much smaller prior austenite grain size and lower hardness observed in the former.
4. 4. 3. Multi-pass weld HAZ.
The double cycle simulation technique enables predictions to be made regarding the HAZ of multi-pass welds where tempering effects will be expected to improve fracture toughness and reduce hardness by reducing dislocation densities and coarsening the carbides. Figs. 7 and 8 and Table 1 show that this is true for regions experiencing a tempering cycle below the upper critical temperature. These regions easily satisfy the NDT requirement at -10 C and have much reduced hardness. However, a further impairment of fracture toughness coupled with hardnesses in excess of
400 HV5 in the grain coarsened HAZ is indicated for regions experiencing a second high peak temperature cycle, with Charpy energy levels reduced to about 34J (25 ft. lb. ) at -10 C, which is also the upper shelf energy level. If such regions exist in a multi-pass weld they would be borderline with respect to the NDT requirement at the minimum service temperature. Such a situation may arise either in regions where two fusion boundaries come into close contact and subsequent weld deposits may then be too far removed to cause effective tempering or by a temper bead positioned too close to a final edge bead HAZ. This may be particularly significant in the final edge bead HAZ which is often a region of appreciable s t r e s s concentration. Oldridge (Ref, 13) has demonstrated the uncertainties involved in accurate positioning of the temper bead,
4 , 4 . 4 . Post-weld heat treatment
Figs. 7 and 8 show that post-weld heat treatment at 650 C for 1 hour is a much more effective and reliable method of improving Charpy performance of the weld HAZ in HY80 than relying on tempering effects occurring in multi-pass welds. Energy absorption is restored quite closely to parent material levels and additional benefit would be derived from s t r e s s -relief. Such structures may be expected to provide complete protection against brittle fracture for the specific application to submarine structures. The possibility of delayed HAZ cold cracking would also be eliminated as is evidenced by the low hardnesses of these structures (Table 2).
There a r e , however, a number of questions which need to be answered before post-weld heat treatment can be considered a s a practical proposition
:-1. Is the strength level of the parent material impaired? Even with a localised post-weld heat treatment it is not possible to avoid reheating parts of the parent material close to the heat treatment temperature and it is essential that the strength level of the parent material is maintained, otherwise the load carrying capacity will be reduced.
2. Are the weld metal properties impaired? Then?- is evidence to suggest that post-weld heat treatment in the temperature range 510 - 650 C may actually impair the notch toughness of the weld metal (Refs. 18 and 31). The overall effect may be a compromise between the benefit derived from reduction in residual s t r e s s e s and the impairment of notch toughness. This will depend on the type of weld metal (i. e. electrode and flux combination).
3. Will the weldment suffer from temper brittleness? There is not much information on this topic although at least one person has reported that HY80 is not susceptible to this form of embrittlement, (Ref. 21). This form of embrittlement is generally associated with slow cooling or holding in the temperature range 620 - 400 C and is a function of chemical composition. Increasing Ni, Cr, and Mn promote susceptibility, while Mo has a counteracting effect.
Until these questions are answered it is not possible to state whether post-weld heat treatment will be beneficial to the post-weldment as a whole. There is an economic penalty to be paid for carrying out post-weld heat treatment so that it is necessary to be able to show substantial benefit before the procedure can be recommended. Heat treatments at other temperatures and times need also to be examined in order to determine the optimum combination and the degree of control which would be required. Accumulated experience with s t r u c t u r e s and p r e s s u r e vessels and information from full scale t e s t s have supported the position that HY80 weldments generally do not require s t r e s s -relieving (Ref. 31). The satisfactory performance of the many welded but n o n - s t r e s s relieved structures and vessels in HY80, stands as undeniable evidence that the steel generally does not require post-weld heat treatment
(stress-relieving). The present work, although confined to a single cast of HY80, has more or less substantiated this view.
5. CONCLUSIONS
1. A marked impairment of fracture toughness, a s measured by the Charpy V notch impact test, and hardnesses in excess of 400 HV5 occur in the grain coarsened and grain refined regions of single-pass weld HAZs in HY80 using a recommended heat input and preheat temperature of 2.1 k J / m m (54 kJ/in) and 120 C respectively. The structures produced are p r e
-dominantly martensitic in c h a r a c t e r and the mechanical properties deteriorate with decreasing distance from the fusion boundary.
2. Little impairment of mechanical properties occurs in regions of the weld HAZ experiencing peak temperatures below the upper critical.
3. Tempering effects occurring in the HAZ of multi-pass welds generally improve Charpy performance and reduce hardness. However, some regions may actually have similar or reduced Charpy performance with hardnesses still in excess of 400 HV5. This is particularly true of the final edge bead HAZ where effective tempering r e l i e s on accurate positioning of the temper-bead.
16
-4. Post-weld heat treatment at 650 C for 1 hour r e s t o r e s Charpy performance of all HAZ structures close to parent material levels and reduces hardness to well below 350 HV5. This is a much more effective and reliable method of improving HAZ properties than relying on tempering effects occurring in multipass welds. Other factors, however, need to be investigated before the treatment can be considered as a practical
proposition.
5. Tentative correlations with NDT temperatures measured by the drop weight or explosion bulge tests and applied to the Pellini fracture - safe design philosophy for submarine structures indicate little risk of brittle fractures initiating from small flaws in non-tempered weld HAZs in HY80. More information, however, is needed to confirm this.
6. HY80 is a suitable replacement for QT35 steel in submarine structures.
R E F E R E N C E S
Steel P l a t e , Alloy, S t r u c t u r a l , High Yield Strength, M i l i t a r y Spec. , M I L - S - 1 6 2 1 6 D (Navy); 12 F e b r u a r y 1959.
Steel P l a t e , Alloy, S t r u c t u r a l , High Yield Strength, M i l i t a r y Spec. , M I L - S - 1 6 2 1 6 E (Navy): 26 J u n e , 1959, a m e n d e d 3 August 1961.
Steel P l a t e , Alloy, S t r u c t u r a l , High Yield Strength, (HY80 and HYIOO) M i l i t a r y S p e c , M I L - S - 1 6 2 1 6 F (Ships), 29 J u n e , 1962.
Steel P l a t e , Alloy, S t r u c t u r a l , High Yield Strength, (HY80 and HYIOO), M i l i t a r y S p e c , MIL S-16216 G (Ships), 27 F e b r u a r y , 1963.
M i n i s t r y of Defence, Ship Dept. , Welding and F a b r i c a t i o n of HY80 Steel, P r o c e s s Spec. No. DG S h i p s / P S / 9 0 2 7 A , J a n u a r y , 1967. M c K e e , A. I. H e l l e r , S. R. F i o r i t i I. , and V a s t a , J. B a k e r , R. G. , W i l k i n s o n , F . , and Newman, R. P . B o n i s z e w s k i , T . . and B a k e r , R. G. Watkinson, F . , B a k e r , R . G . , and T r e m l e t t , H. F . N i p p e s , E . F . Savage, W. F . Allio, R . J . Dolby, R. E . O l d r i d g e , D. and R e c e n t s u b m a r i n e d e s i g n p r a c t i c e s and p r o b l e m s . T r a n s . Soc. Nav. A r c h i t . M a r . E n g r s . N. Y. 67, p . 632, 1959. An e v a l u a t i o n of HY80 s t e e l a s a s t r u c t u r a l m a t e r i a l for s u b m a r i n e s , Naval E n g . J n l . p. 29, F e b r u a r y , 1965. p. 193, A p r i l , 1965. The m e t a l l u r g i c a l i m p l i c a t i o n s of
welding p r a c t i c e a s r e l a t e d to low alloy s t e e l s ; Second C o m m o n w e a l t h Welding C o n f e r e n c e , p. 125, 1965.
Heat affected zone cold c r a c k i n g in low alloy s t e e l s ; Second C o m m o n w e a l t h Welding C o n f e r e n c e , p. 117, 1965.
Some thoughts on the p r o b l e m of welding high t e n s i l e s t e e l s ; Welding in Shipbuilding Symposium, I n s t i t u t e of Welding, p. 3 1 , 1962.
Studies of the weld heat affected zone of T - 1 Steel, Weld. J n l . . 36 (12), p . 531-s, 1957.
The effects of welding on the f r a c t u r e t o u g h n e s s of quenched and t e m p e r e d low alloy s t r u c t u r a l s t e e l s . Weld. Inst. Rep. C 2 0 1 / 1 0 / 6 8 , 1968.
The effect of a t e m p e r i n g bead on the p r o p e r t i e s and s t r u c t u r e of the weld h e a t affected zone in Q T 3 5 s t e e l ,
D . A . E . 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 , 1964.
18
14. Clifton, T . E . . and George M. J.
Design and construction of a weld heat affected zone simulator. Met. Constr. ,
1 (9), p 427, 1969. 15. 16 Coward, M. D. Smith, E, 17 Kellock, G. T. B. 18. Dolby, R. E. 19. Inagaki, M. , Uta, M. and Wada, T.
The structure and properties of the heat affected zones of welds in mild and low alloy steels, M.Phil. Thesis, University of London, 1967.
Unpublished work, Cranfield Institute of Technology, Cranfield.
A study of simulated weld heat affected zone structures and properties of HY80 low alloy steel, D.A. E. Thesis, Cranfield Institute of Technology, Cranfield, 1969. Private communication.
A new apparatus for determining SH-CCT diagram for welding and its application to high strength steel. Trans. Nat. Res. Inst, for Metals, 6 (6) p. 386, 1964. 20. 2 1 . Puzak, P. P. , Eschbacher, E. W., and Pellini, W. S. Puzak, P. P. , and Pellini, W. S.
Initiation and propagation of brittle fracture in structural steels; Weld. J n l . , 31, (12), p. 561-2, 1952.
Evaluation of the significance of Charpy Tests for quenched and tempered steels, Weld. J n l . , 35 (6), p. 275-s, 1956. 22. Pellini, W . S . . and
Srawley, J. E.
Procedure for the evaluation of fracture toughness of p r e s s u r e - v e s s e l materials, N. R . L . Report 5609, Washington D. C. June, 1961. 23. Pellini, W. S. and Puzak, P. P . 24. Pellini, W . S . , 25. Winn, W.H. 26. Smith, E.
Fracture analysis diagram procedures for the fracture-safe engineering design of steel structures, N. R. L. Report 5920, Washington D. C. March, 1963.
Evaluation of engineering principles for fracture-safe design of steel structures, N. R . L . Report, 6957, September, 1969. Welding higher strength steels in war-ship construction. Second Commonwealth Welding Conference, p. 209, 1965.
Unpublished work, Cranfield Institute of Technology, Cranfield.
Smith, E . , and A p p s , R. L. G r o t k e , G. E . , W e s s e l , E . T. , and H a y s , L . E . Savage, W. F . , and O w c z a r s k i , W. A . , Doty, W . D . D o r s c h u , K. E . , and L e s n e w i c h , A.
Effect of welding and p o s t - w e l d heat t r e a t m e n t on Q T 3 5 , Cranfield R e p o r t Mat. No. 2, Cranfield I n s t i t u t e of Technology, Cranfield, 1970. W e l d a b i l i t y and heat affected zone t o u g h n e s s of HY150 s t e e l . W e l d . J n l . 43 (6), p. 2 6 5 - s , 1964.
D e v e l o p m e n t of a h i g h - s t r e n g t h , tough, w e l d a b l e , s t r u c t u r a l s t e e l . W e l d . J n l . , 43 (5), p . 2 1 5 - s , 1964.
The m i c r o s t r u c t u r e and notch i m p a c t b e h a v i o u r of a welded s t r u c t u r a l s t e e l , W e l d . J n l . , _45, (2), p. 5 5 - s , 1966.
Welding of quenched and t e m p e r e d s t e e l s , Weld. J n l , 30, (7), p. 2 8 9 - s , 1965. D e v e l o p m e n t of a f i l l e r m e t a l for a high t o u g h n e s s alloy plate s t e e l with a
m i n i m u m yield s t r e n g t h of 140 K s i . Weld. J n l . , 4 3 , (12), p. 5 6 4 - s , 1964. K i l p a t r i c k , I. P r i v a t e c o m m u n i c a t i o n .
M a t e r i a l Condition 1st Cycle 1275 1275 1275 1275 1275 930 930 930 765 765 P a r e n t 2nd Cycle 1275 930 765 650 765 650 650 T r a n s v e r s e Longitudinal T e n s i l e P r o p e r t i e s 0. 2% P r o o f S t r e s s N / m m 2 (tonf/in2 ) 939 (61.0) 1000 (65.1) 1028 (66.9) 793 (51.5) 930 (60.4) 910 (59.1) 675 (43.9) 771 (50.1) 602 (39. 1) 577 (37.5) 576 (37.4) 594 (38.6) U T S N / m m 2 (tonf/in2) 1328 (86.3) 1273,'-- (82.6) 1350 (87.7) 1023 (66.4) 1048 (68.1) 1080 (70.1) 884 (57.4) 881 (57.2) 755 (49.0) 720 (46.8) 719 (46.7) 735 (47.8) Reduction of A r e a % 51 53 58 57 61 62 62 66 72 71 72 70 Mean H a r d n e s s HV5 423 412 43 7 308 352 353 318 1 279 1 278 234 228 230
1275 1275 1275 1275 1275 930 930 930 765 765 _ 1275 930 765 650 -765 650 -650 735 (47.8) 751 (48.8) 770 (50.0) 705 (45.8) 743 (48.3) 608 (39.5) 594 (38.6) 607 (39.4) 565 (36. 7) 556 (36. 1) 837 (54.4) 827 (53.7) 825 (53.6) 813 (52.8) 815 (53.0) 706 (45. 9) 698 (45.3) 705 (45. 8) 705 (45.8) 680 (44.1) % 70 72 63 70 68 74 73 72 78 73 268 272 273 255 273 231 238 232 230 228 T A B L E 2 M e c h a n i c a l p r o p e r t y d a t a for s i m u l a t e d s p e c i m e n s , p o s t c y c l e heat t r e a t e d at 650°C.
o UJ ft:
I
UI O. 1300 1200 1100 1000 900 800 700 600 500 400 300 200 100\-( Q ) , \-( c } : MEASURED THERMAL CYCLES (b),(d):COMPUTED THERMAL CYCLES
10 20 30 40 50 60 70
TIME , SEC.
FIGURE.I. THERMAL CYCLES PRODUCED IN THE PARENT PLATE ADJACENT TO THE WELD ,((a) AND (c)) , AND THOSE PRODUCED BY COMPUTATIONAL TECHNIQUES , ( ( b ) AND (d)).
(b) carbon extraction replica
FIG. 2 PARENT PLATE MICROSTRUCTURE.
M (O UI z a K K UJ > 480 UO 400 360 3 2 0 280 -240 200 1 ' ' > 1 f
h •' ^
r \ < L \ ' 1 * X L w" ^ 1 —1—r .1 i i V 1 \ » • 1 I 1 FUSION HAZ ' 1 NI '• . 1 . • • 1 1 1 • 1 • 1 • 1 • I BOUNDARY —^^ BOUNDARY VlW SrM^l
WELD B E N D - X ^ ^ ^ O ^ , ] PARENT PLATE—"'-•^'^-l--^''^ JUNE OF HARDNESS SURVEY '
A ;• HVO-5 Ac 1 . 1 . > l . - L .
-]
\ j4
Ac, J . ! _ . 1 1 1 0-4 0 8 1-2 1-6 2 0 2-4 DISTANCE FROM FUSION BOUNDARY, mm2 8 3-2
FIG, 4 THE NATURE OF BANDING IN THE PARENT PLATE AFTER THERMAL
ENERGY ABSORBED , f t I b t . X -n m 5 > c •^s -• V ^ m > m X "" m m z o
-• -i
j^ X o o -n X -• d • — X m2?
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FIG. 15 POST-CYCLE HEAT TREATED STRUCTURES OF SINGLE AND DOUBLE CYCLED SPECIMENS WITH INITIAL CYCLE TO A PEAK TEMPERATURE OF 930°C
