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Anewseriesofadvanced3Cr-Mo-Nisteelsforthicksectionpressurevesselsinhightemperatureandpressurehydrogenservice

ARTICLEinJOURNALOFMATERIALSFORENERGYSYSTEMS·DECEMBER1984

DOI:10.1007/BF02833446

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A New Series of Advanced 3Cr-Mo-Ni Steels for Thick Section Pressure Vessels in High

Temperature and Pressure Hydrogen Service

R . O . R I T C H I E , E. R. P A R K E R , P. N. S P E N C E R , and J. A. T O D D

A new series of 3Cr-Mo-Ni steels has been developed for use in thick section pressure vessels, specifically for coal conversion (high temperature and high pressure hydrogen) service. The new steels rely on minor alloy modifications to commercial 2.25Cr-lMo (ASTM A387 Grade 22 Class 2) steel. Based on evaluations in relatively small heats (55 kg), the experimental alloys, which employ additions of Cr, Ni, Mo, and V, with Mn at 0.5 wt pct and C at 0.15 wt pct, display improved properties compared to commercial steels. Specifically, they show significantly improved hardenability (i.e., fully bainitic microstructures following normalizing of 400-mm (16-in.) plates), enhanced strength (i.e., yield strengths exceeding 600 MPa), far superior hydrogen attack resistance and better Charpy V-notch impact toughness, with comparable tensile ductility, creep rupture resistance and temper embrittlement resistance. The microstructural features contributing to these improved mechanical properties are briefly discussed.

INTRODUCTION

The development of second and third generation coal con- version systems, such as proposed large-scale coal liq- uefaction and gasification processes, has necessitated the design of large thick-section pressure vessels. L2.3 Materials

requirements for such vessels include weldable steels with sufficient hardenability to maintain good mechanical prop- erties throughout plate sections up to a maximum of roughly 400 mm (16 in.). In addition, yield strengths in excess of 350 MPa are required with sufficient toughness, creep rup-

ture, fatigue and environmental degredation resistance to withstand mechanically and environmentally hostile envi- ronments, which in certain instances involve hydrogen plus hydrogen sulfide atmospheres at temperatures of - 5 5 0 ~ and hydrogen gas pressures up to 20 MPa (3000 psi). 3-6

R. O. RITCHIE, E. R. PARKER, and P. N. SPENCER are Pro- fessor, Professor Emeritus, and Research Engineer, respectively, in the Department of Materials Science and Mineral Engineering, University of California, Berkeley, CA 94720. J.A. TODD is Assistant Professor, Department of Materials Science, University of Southern California, Los Angeles, CA 90089. Presented at American Society for Metals WESTEC '84, March 1984, Los Angeles, CA.

J. MATERIALS FOR ENERGY SYSTEMS �9 1984 AMERICAN

Historically, the material favored for hydrogen service pressure vessel construction has been 2.25Cr- lMo steel, 6,7 such as the normalized and tempered ASTM A 387 Grade

22 Class 2. Although this steel in general has ideal mechani- cal properties, it does not have sufficient hardenability to produce the fully bainitic microstructures necessary to pro- vide the desired low and elevated temperature properties in normalized plate sections of the required 300 to 400 mm (12 to 16 in.) thicknesses. 8 Moreover, there is also some doubt as to the resistance of 2.25Cr-1Mo steel to hydrogen attack under the most severe in-service conditions, 9-13 where elevated temperature and pressure gaseous hydrogen envi- ronments can result in an internal react ion be tween ingressed hydrogen and carbides in the steel. This leads to decarburization, cavitation and, in extreme cases, fissuring from the formation and growth of methane bubbles at inter- faces such as grain boundaries. ~4'~5'16

Both laboratory research and in-service exper ience have shown that additions of certain alloying elements, specifically carbide stabilizing elements such as vanadium, niobium, and particularly chromium and molybdenum, can have a markedly beneficial effect on hydrogen attack resistance through the precipitation of stable carbides.~7"~8

However, there is some concern over the elevated temper-

SOCIETY FOR METALS VOL. 6, NO. 3, DECEMBER 1984 151

ature strength of such modified steels, due to their greater tendency to precipitate M23C6 carbides which tend to coarsen at service temperatures. 19.20 Furthermore, commer- cial 3Cr-lMo compositions, such as ASTM A 387 Grade 21 Class 2, also have insufficient hardenability for section sizes in excess of about 200 mm. 2~'22

Over the past few years, several major alloy design pro- grams have been instigated to develop superior alternatives to commercially available heavy section 2.25Cr-lMo pres- sure vessel steel for elevated temperature hydrogen ser- vice. 5"39-25 Notable amongst these studies are the Japanese

Steel Works 3Cr-lMo steel microalloyed with Ti, V and B 24 and the Climax Molybdenum Company 3Cr-l .5Mo steels containing 0.1 pct V and 1 to 1.4 pct Mn with 0.12 pct C maximum. 19,21,23

In the current alloy development program the objective was to design an improved hydrogen attack resistant steel of higher strength to permit applications with thinner section sizes, thereby saving in weight and cost. Similar additions of Mo and Cr were employed to improve elevated tem- perature strength, oxidation resistance and resistance to hy- drogen damage. However, 0.5 to 1 pct Ni was also added for hardenability coupled with 0.2 pct V for creep resistance and grain refinement. In particular, the Mn content was limited to a nominal 0.5 pct in order to minimize potential problems, from temper embrittlement susceptibility, 19,26,27,28

excess retained austenite following slow cooling* and hand-

*Both banding and the transformation of excess retained austenite on tempering can lead to unexpected local susceptibility to hydrogen attack due to non-uniform distributions of unstable alloy carbides. ~2

ing, which can be promoted by higher Mn contents. The specific compositions (Table I), consisting of two 3Cr-1Mo- 1Ni steels, with and without 0.2 pct V (termed Steels C and B, respectively), and two 3Cr- 1.5Mo-0.5Ni steels, with and without 0.2 pct V (termed Steels E and D, respectively), are compared with a conventional 2.25Cr-lMo steel of similar purity and steelmaking practice (termed Steel A).

Such modified 3Cr-Mo-Ni steels are found to be far superior to commercial 2.25Cr-lMo steels and to compare very favorably with other advanced 3Cr-Mo materials. The microstructural features contributing to the improved hydro- gen attack resistance, as well as the superior hardenability strength and toughness, and comparable ductility, temper

embrittlement resistance and creep rupture properties, are discussed in this paper.

EXPERIMENTAL PROCEDURES

The four experimental 3Cr-Mo steels and the reference 2.25Cr-lMo steel were produced as 55 kg (125 lb) labora- tory heats by Climax Molybdenum Company using vacuum induction melting and casting under argon atmospheres. Chemical compositions in wt pct are listed in Table I. The ingots were subsequently upset-forged and cross-rolled to approximately 30 mm thick plates before austenitizing for 1 hour at 1000 ~ Following austenitization, samples were either quenched into agitated oil (i.e., cooling rate

1200 ~ or subjected to slow continuous cooling, at 8 ~ through the transformation range, using a pro- grammable induction furnace. Based on cooling rate data supplied by Lukens Steel Company, these treatments simu- lated both the surface and quarter thickness (0.25 T) loca- tions, respectively, of a 400-mm (16-in.) thick plate during accelerated surface cooling (herein referred to as nor- malizing). Tempering was carried out in neutral molten salt baths at temperatures of 650 ~ and principally 700 ~ for a range of times varying from 1 to 1000 hours. In terms of the tempering parameter, P, sometimes referred to as the Larson Miller parameter and defined as T[20+log t] • 10 -3 where T and t are the tempering temperature (in Kelvin) and time (in hours), respectively, these treatments

represent a variation in P between 19.0 and 22.4 (or be- tween 34.2 and 40.3 for temperatures in Rankin).

Room temperature uniaxial tensile tests were conducted on 6.4-mm diameter specimens of 32-mm gauge length, machined in the longitudinal direction of the plate, accord- ing to ASTM Standard E8-69. A displacement rate of 0.5 mm/min was employed. Standard Charpy V-notch impact specimens were also prepared in the longitudinal direction, and tested according to ASTM Standard E23-72 over a temperature range from - 1 9 6 ~ to 160 ~ Constant load creep-rupture tests were performed in air using smooth round bar longitudinal specimens of initial diameter 6.4 mm and gauge length 25 ram, according to ASTM Standard E-139. Tests were carried out at 560 ~ with engineering stresses between 138 and 345 MPa (20 to 50 ksi).

Susceptibility to temper embrittlement was evaluated by testing a further series of Charpy impact toughness speci-

Table I. Chemical Composition, in wt pct, of Reference 2.25Cr-1Mo and Experimental 3Cr-Mo Steels Tested

Designation C Mn Si Cr Ni Mo V P S AI

Steel A 0.15 0.48 0.21 2.24 - - 1.01 - - 0.006 0.0034 0.014 Steel B 0.15 0.50 0.23 3.00 0.98 1.03 - - 0.008 0.0040 0.010 Steel C 0.14 0.49 0.24 2.99 0.98 1.02 0.21 0.006 0.0039 0.009 Steel D 0.15 0.47 0.22 3.00 0.50 1.50 - - 0.006 0.0036 0.008 Steel E 0.14 0.49 0.22 2.95 0.50 1.50 0.21 0.006 0.0034 0.014

152 VOL. 6, NO. 3, DECEMBER 1984 J. MATERIALS FOR ENERGY SYSTEMS

mens which had been step-cooled, rather than air cooled, following tempering at 700 ~ The step-cooling treatment 26 involved holding for progressively longer times at progres- sively decreasing temperatures from 595 ~ to 470 ~ as illustrated schematically in Figure 1.

Susceptibility to hydrogen attack was evaluated with re- spect to both uniaxial tensile and impact toughness proper- ties by prolonged exposure of oversize specimen blanks to gaseous hydrogen atmospheres at temperatures of 550 ~ to 600 ~ and at pressures of 14 to 18 MPa (2000 to 2800 psi) for times up to 1000 hours, prior to machining and testing.

Microstructures were examined with both optical and electron microscopy. Thin foils for transmission electron microscopy (TEM) were prepared from 0.6 mm steel slices by chemically thinning to 0.05 mm in a hydrofluoric acid/hydrogen peroxide solution before electropolishing at room temperature in chromium trioxide/acetic acid solution. Foils were examined using Phillips EM301 and 400 STEM electron microscopes at 100 kV. Further analysis of the carbide compositions was performed on the scanning transmission electron microscope (STEM) using extraction replicas. Quantitative estimates of the percentage of retained austenite were assessed using both X-ray and magnetic saturation techniques.

RESULTS

Strength and Ductility

The room temperature uniaxial tensile properties of the four 3Cr-Mo-Ni steels, after slow-cooling (8 ~ and tempering at 700 ~ for various times between 1 and 1000 hours, are shown in Figure 2. These heat-treatments represent a range of tempering parameters (temperatures in Kelvin) between 19.5 and 22.4. It is apparent that all four

UNEMBRITTLED (Slow cooled and tempered)

IOOO*C, I hr

700Oc, 4 h r

~eOC/min Air COOl

EMBRITTLED (Slow cooled, tempered and step cooled)

1000~ I hr

~ 700eC, 4 hr s~*c,, h,

/ \ F-~ ~ o ' c , ,5., I \ / \ 495"C, 48 hr

I \ / Furnace cool ~ " ~ SI . . . . . I "~ '~ STEP COOL to 315~

~8~ Air cool Air COOl

Fig. 1 - Schematic representation of the heat treatments used to assess susceptibility to temper embrittlement.

J. MATERIALS FOR ENERGY SYSTEMS

1400

120(

Temperin~ Parameter P = T (K) [20-1- Io9 t(hrs)J=lO -3

19 20 21 22 1 I I I

- HEAT-TREATED

~ ~O- -~ ~ (Y" . . . . . .

o o ~ ~ ~'*/o Red'n. Area

80

~0

A s ~E

al t/)

600

400

200

r\ ~ ,.,@,,. ItS

--%. )~ x A

. , o o

yield Strengt h '~ '~ "~...~ ? . ~ . ' - - z,.."~ ~ t " - ' ' ~ . . . . i l l

- ' : l " - - -

~" O E ~ ' " " 7 " " ~ O [] 0 R.A.

-- % Elont.jotion 4L ~ ~ (~ UTS

& (ll rl ~ YS . . . . . �9 �9 �9 �9 Elon(J. ~ . . . .

o - - f [ I J I o 0 I I 0 I00 1000

Tempering Time (hr$) of 7OO"C

Fig. 2 - - R o o m temperature uniaxial tensile properties for experimental 3Cr-Mo-Ni steels, slow-cooled and tempered at 700 ~ as a function of tempering time and tempering parameter (in Kelvin).

steels show similar mechanical properties (the V-containing C and E steels are marginally stronger), with ductilities fairly constant at pet RA between 70 and 80 pet, and with tensile strengths above 600 MPa (87 ksi), even at the longest tempering times.

Toughness

The Charpy V-notch impact toughness, as a function of test temperature for the four 3Cr-Mo-Ni steels, slow-cooled (8 ~ and tempered 4 hours at 700 ~ (P = 20), is shown in Figure 3. The toughness of Steels C, D, and E

1 5 0 - SC 8tT ot 700~ 4 h r s l O Steel e ~ . - ~ 200 / , , Stee, C /

os,e.,o o / / / / ~ | o s, ee~ E ]~50 ~,

'~176 / 7/- -

st., 8 E

l fst"'c//7--s'"'~ ],oo ,5o -

~176 // 1 o

0 0 I 1 I I l 1 [ 0 - 2 0 0 -I 0 0 0 100 2 0 0

Test Temperoture (~ Fig. 3 - - C h a r p y V-notch impact toughness transition curves for experi- mental 3Cr-Mo-Ni steels, slow-cooled and tempered (SC & T) at 700 ~ f o r 4 h.

VOL. 6, NO. 3. DECEMBER 1984 153

is clearly similar with a 40 ft-lb (equivalent to 54 J) ductile/brittle transition temperature just below 0 ~ and an upper shelf energy above 176 J (130 ft-lb). The 3Cr- 1Mo-lNi Steel B, however, is distinctly tougher with a transition temperature be low-100 ~ and an upper shelf energy above 200 J (150 ft-lb). 2~ This steel shows similar high toughness when oil quenched and when tempered 4 hours at 650 ~ (P -- 19), as seen by the (unexposed) data in Figure 4.29

Temper Embrittlement Resistance

Charpy V-notch toughness transition curves for the four 3Cr-Mo-Ni steels in a temper embrittled condition, induced by a step-cooling treatment (Figure 1) following tempering 4 hours at 700 ~ are compared in Figure 5 with transition curves of corresponding unembrittled microstructures. With embrittlement, all steels show a reduction in upper shelf energy of between 10 and 40 J and a shift (AT) between 15 ~ to 55 ~ in the 40 ft-lb transition temperature to higher temperatures (Table II). Although behavior is again some- what similar in the four steels, the 3Cr- lMo-lNi Steel B still displays the lowest ductile/brittle transition temperature.

200 I I I I [ I I

3 C t - I M o - I N i S T E E L B 2 5 0

z ~ / o

>" b .V SC+ T?OO~C

,~....-,--"T . . . . . . . n. ~50

U N E X P O S E D / l 7 , ' , 0 7 / ~ / / , '.E'

- / P / , ; / , ' / ' . , o o , , ~ o - ~ . . ~ ~o / / / / f. �9 0Q~'TTOO'C+HE

0 ~ , ~ r . ~ ~ - . r 0 ~ - ] [ ~ I I I I - 2 0 0 -I00 0 r 0 0 2 0 0

T E M P E R A T U R E (*C)

Fig. 4 - - Charpy V-notch impact toughness for experimental 3Cr- l Mo- l Ni Steel B, oil quenched (OQ) or slow-cooled (SC) and tempered (T) 4 h at 650 ~ or 700 ~ in the prior hydrogen exposed (HE) condition (600 ~ 17 MPa pressure) and unexposed condition. ~9

Hydrogen Attack Resistance

In Figure 6, the room temperature uniaxial tensile data for the four 3Cr-Mo-Ni steels are given, as a function of tem- pering time at 700 ~ following slow cooling from 1000 ~ for samples which have been previously exposed for 1000 hours to 14 MPa (2000 psi) hydrogen gas at 550 ~ By comparing with the scatter bands, which represent the

ISC

t d - IOC

E

~ 50

2 0 -200

I I I I I I I I

Steel B ( 3 C t - I M o - I N i ) ZOO SCS=TOt700C 4hr

/ ~ ~ T e m p e r Embrittled

""-- I I I O I I I I I 0 -I00 0 I OO 200

Test Temperoture (~

150 I

"' I00

~ -

~ 50--

U ~

-2OO

I I I f I I f |

Steel C ( 3 C r - I M o - I N i - 0 , 2 V )

SC 8= Tot 700 ~C; 4 hrs

Unembrittled ~

l / "Temper Embrittled J / (Step cooled)

l ! I~t-~1-- I I I I I -IOO 0 I00 200

Test Temperoture (~C:)

200

150

w i i

IOO

50

A

-'12 1 5 0 -

f -

w

~ 1oo

E

o 50

o ~J O

-2OO

! ! I I i

Steel O ( 3 Cr- I.S Mo -0.,'5 Ni ) SC & T ot 7OOeC 4 hrs ~ ' ' " ~

Unembrittled ~ / Y ~

/ / ~Temper Embrittled J (Step cooled)

I I~"- I I 'r'l I I I I -IOO 0 I0o 200

Test Temperoture (='C)

I ! !

- 200

-- 150

50

O

A

i $ 0 -

e~ c LIJ - 100

= 5 0 - -

r

-200

I I I ! I ! ! I

Steel E (3Cr-I.SMo-O.SNi-O.ZV)

SCS=Tot 7OO~C 4hrs S

/ / "~t-emper Emlxittled

I I V " - I I I I I I

-I0O 0 IOO 200 Test Temperoture (*C)

200

150

ioo o

50

0

Fig. 5--Charpy V-notch impact toughness transition curves for experimental 3CR-Mo-Ni steels, slow-cooled and tempered (SC & T) 4 h at 700 ~ in the temper embrittled (step-cooled) and unembrittled conditions.

154 VOL. 6, NO. 3, DECEMBER 1984 J. MATERIALS FOR ENERGY SYSTEMS

Table II. Effect of Temper Embrittlement on Charpy V-Notch Impact Toughness

40 ft-lb (54 J) Transition Temp. Upper Shelf Energy Shift

Steel* Unembrittled Temper Embrittled Unembrittled Temper Embrittled AT

(J) (J) A - 13 ~ - - 122 - - - - B - 1 0 5 ~ - 4 0 ~ 206 165 65 ~ C - 25 ~ - I 0 ~ 186 175 15 ~ D - 20 ~ 25 ~ 203 175 45 ~ E - 10 ~ 15 ~ 176 163 25 ~

,o1, ~ered ,. hr

Temper ing P a r a m e t e r 1501 I I I I - P = T(K) [ :>0+ log t (hrs)] x 10 -3 I ,,~ ~o ~,, ~,~ ~ I H~'T-TREATED S'EELS

. . . . . . . . . . . . . ' =

'x " ,=t O ~ ~ _ " . . . ~ " I I - . e _ ~

" " ~ " - L ~ . " - . - ' < ~oo �9 ~ : ~

Yield Strength / ~ ~

- - :>-~- : : - - . ,_" ~ ~EM.ER,.G T,ME ,..S,"OOAT ,OO'~ 't~176176 ( a )

i ~ . _ ~ . - e - - ~ :o t5o , i , , - .~ /~ , . . .~ . . . . . .~ .~" t C B O E

*Slow cooled (8 ~ tempered 4 hr at 700 ~

2 0 0

150 i

I 0 0 ~

5O

:0

2 0 0

1 4 0 0 [ - HYDIRoGE__ N EXPOSED I 80 I \

I ~ --'" 0 0 0 0

I %'" o _ . . , a . . . _ _ o . _ . L ~ _ . . . . . . . . , ~ o o ~ / ~ . . . . . . . o ~ ~...d'n~..~

_ I ' , \ L > < 2 ' s -

A 0 0 0 R.A. 20C "&Elongation • ~ ~ $ UTS

& (! I ! ~, YS . . . . . �9 �9 �9 �9 Elong.-- . . . .

o ~ ' f = I = o I IO I 0 0 I 0 0 0 T e m p e r i n g Time (hrs) at 7 0 0 " C

Fig. 6 - - R o o m temperature uniaxial tensile properties for prior hydrogen exposed 3Cr-Mo-Ni steels, slow-cooled and tempered at 700 ~ as a function of tempering time and tempering parameter. Hydrogen expo- sure for 1000 hr at 550 ~ with 14 MPa pressure. Scatter bands repre- sent bounds of data for corresponding unexposed conditions, shown in Figure 2.

._1 O- h

) -

IOC Z I,u

U <I O. ~E

o. 5C r r

"I" t.3

bounds of the corresponding data on unexposed samples from Figure 2, it can be seen that, apart from a slight decrease in pet RA in the higher strength conditions, there is little evidence of hydrogen attack damage in these steels (tempered at 700 ~ for this hydrogen exposure.

A similar result is obtained when comparing the effect of hydrogen attack damage on room temperature Charpy V- notch toughness (F'gure 7). In Figure 7(a), the toughness of unexposed samples is plotted as a function of tempering

S?EL B (HEAT-TREATED) 7 " I / / . . . . . . . . . / l-

,, 7/- , - , ~ 0

/ / f

I I ! I I

I 0 / / 0 / /

0 ~.(I 0

/ . . . . 0 - - - - I

f o o ~ ~ , , / 0 t

/ / O / 0 / -

/ !

/ O /1 ~ N - E X P O S E D

/ AND / HEAT TREATE'D I I

-E O STEEL B -

~ D STEEL O (HEAT-TREATED) O STEEL E

I I I I0 I00 I000

TEMPERING TIME (HRS) AT 700=C

150

7, tlJ _J

IOO~ o

5 0

(b) Fig. 7 - - R o o m temperature Charpy V-notch impact energy as a function of tempering time at 700 ~ for slow-cooled 3Cr-Mo-Ni steels in the (a) heat-treated (unexposed) and (b) hydrogen-exposed conditions. Hydrogen exposure for 1000 h at 550 ~ with 14 MPa pressure. Scatter bands repre- sent bounds of data for unexposed conditions.

J. MATERIALS FOR ENERGY SYSTEMS VOL. 6, NO. 3, DECEMBER 1984 155

time at 700 ~ following slow-cooling from 1000 ~ As noted above, Steel B is distinctly tougher and reaches a high room temperature toughness, corresponding to upper shelf ductile fracture by microvoid coalescence, after only short tempering times at 700 ~ The other steels show mixed mode transitional behavior, with consequent lower tough- ness, as their respective transition temperatures are closer to room temperature ( c . f . , Table II). Following prior 1000-hour exposure to 14 MPa hydrogen at 550 ~ the toughness in all steels is essentially unchanged [Figure 7(b)], with no apparent differences in fracture morphology compared to corresponding unexposed conditions. Steel B, however, does show a slight decrease in toughness. This may be more associated with entrapped hydrogen rather than hydrogen attack damage from internal voids and fissures, because subsequent re-heating of prior damaged samples for 80 hours at 275 ~ (which bakes out dissolved hydrogen) returns the toughness to undamaged levels ( i .e . , excess of 135 j).25

The effect of more severe hydrogen exposures [corre- sponding to 17 to 18 MPa (2500 to 2800 psi) pressures] is shown in Table III and Figure 4 for the 3Cr-lMo-lNi Steel B, with samples tempered either at 650 ~ or 700 ~ (4 hours) following either oil quenching or slow-cooling (8 ~ z9 For all structures, some small degree of soft- ening can be seen (tensile strengths reduced by 7 to 17 pet), together with a small decrease in upper shelf toughness (by between 2 and 22 pct) and a shift in the ductile/brittle transition temperature (by 65 ~ to 97 ~ Fracture surfaces on broken Charpy specimens were predominately inter- granular close to the notch in prior hydrogen exposed sam- pies, but there was no evidence of voids on the intergranular facets which are so characteristic of hydrogen attack dam- aged structures. ,2

By comparison with Figure 5, it is felt that this decrease in toughness following prolonged hydrogen exposure is again probably not predominately due to hydrogen attack, but rather to temper embrittlement occurring during the ther-

1 I I I I [ I I

,501 S,~e___L~._(3Cr-' Mo-1 N~______) ~ ZOO Slow cooled ond tempered f

at 700 *C/4 h

<u - 150

~J~ IOC Unembrittled\ �9 �9

o. ~ , ~ / ~ ~--Temper Embrittled '~

/ C/~Hydrogen Exposed 50 ~ ~ 1 1 Z {600eC/IOOOh/17MPo) o

J= u J I I l I 0

-200 - I00 0 I O0 200 Test Temperature (*C)

Fig. 8 - - E f f e c t o f prior temper embrittlemcnt and hydrogen attack on Charpy V-notch impact toughness of experimental 3Cr-1Mo-1Ni Steel B, slow cooled and tempered 4 h at 700 ~

mal exposure. 25 In fact, duplicate Charpy specimens, slow cooled and tempered 4 hours at 700 ~ when heated for 1000 hours at 550 ~ in an inert atmosphere, were found to have similar toughnesses to the hydrogen exposed samples. Furthermore, the toughness of step-cooled (temper em- brittled) samples in this steel was almost identical to that following hydrogen exposure (Figure 8).

Thus it is apparent that there is only a small influence of hydrogen attack damage in these steels, with the sus- ceptibility to embrittlement being far less than that reported for 2.25Cr-1Mo steel or for 3Cr-1Mo-1Ni steels containing higher ( i .e . , 1 pet) Mn contents. ]2

Creep Rupture Resistance

Creep-rupture data for the four 3Cr-Mo-Ni steels (slow- cooled and tempered 4 hours at 700 ~ are shown in Fig- ures 9 and 10. Results at 560 ~ are compared to 2.25Cr-lMo steel :~176 in terms of applied stress vs life (Figure 9) and creep rupture ductility (pet RA) vs life (Figure 10). On the basis of these data, the elevated tem- perature properties of the 3Cr-Mo-Ni steels are clearly

Table lU. Effect of Hydrogen A t t a c k on Uniaxial Tensile and Toughness Properties in 3 C r - l M o - l N i Steel B

Upper Shelf 40 ft-lb (54 J) Yield Stress U.T.S. pct Redn. Charpy Energy Transition Temp.

Condition (MPa) (MPa) pet Elong.* Area (J) (~

UE HE UE HE UE HE UE HE UE HE UE HE

Slow-cooled, 4 hr at 650 ~ 691 552 781 647 Slow-cooled, 4 hr at 700 ~ 627 552 779 661 Oil Quenched, 4 hr at 650 ~ 660 537 782 638 Oil Quenched, 4 hr at 700 ~ 545 502 664 616

20 22 75 73 189 162 - 1 0 0 ~ - 2 3 ~ 24 23 73 73 206 161 - 1 0 5 ~ - 4 0 ~ 24 22 76 73 190 186 - 1 1 5 ~ - 5 0 ~ 24 25 76 71 230 178 - 1 2 0 ~ - 2 3 ~

*On 32 mm gauge length. Legend: UE = Unexposed, HE = Hydrogen Exposed (600 ~ 17 MPa)

156 VOL. 6, NO. 3, DECEMBER 1984 J. MATERIALS FOR ENERGY SYSTEMS

Fig. 9--Creep rupture properties at 560 ~ of experimental 3Cr-Mo-Ni steels, slow-cooled and tempered 4 h at 700 ~ Results are compared with previous data 2~176 on 2.25Cr-lMo steel.

I 0 0

*~ 8 O

6 O

"o

I I I I

-- -- 0 . . . . . (]5 0 �9 0 �9 C~ r

00 13 �9 ~ O

SC a T o t 700~ 4 hrs; Tests ot 560 "C

1 I

�9 Steel A (Bose 2 .25 Cr - IMo) 0 Steel B ( 3 C r - I M o - I N i )

Steel C ( 3 C r - I Mo-I Ni-O.?V) I~ Steel D ( 3 C r - I 5 Mo-OSNi ) 0 Steel E (3Cr -I .5 MO-0.5 Ni-O.2V)

I I I I0

Rupture Time (hours) 0.I I 00 I 0 0 0 I0,000

Fig. 101Creep rupture ductility (pct RA) v s time to rupture for experi- mental 3Cr-Mo-Ni steels, slow-cooled and tempered 4 h at 700 ~ from constant load tests at 560 ~ at stresses between 138 and 345 MPa.

comparable with those of 2 .25Cr- lMo steel. Furthermore, at longer lives (above 1000 hours), the 3Cr-Mo-Ni steels actually out perform commercial 2.25Cr-lMo alloys, par- ticularly in the V-containing C and E steels which tend to display the best performance at the lower applied stresses.

Hardenability

For cooling rates as slow as 8 ~ typical of the condi- tions at quarter thickness of a 400-mm thick plate during normalizing, commercial 2 .25Cr-lMo steel generally has inadequate hardenability to ensure fully bainitic micro- structures, as evidenced by the substantial proportion ( - 4 0 pct) of the polygonal ferrite in as-cooled structures [Figure ll(a)]. The experimental 3Cr-Mo-Ni steels, con- versely, with their increased Cr + Ni additions, show vastly enhanced hardenability and 100 pct granular bainitic structures after a similar slow cooling rate of 8 ~

[Figure 1 l(b)].

Microstructures

As noted above, the microstructure of all four 3Cr-Mo-Ni steels was found to be 100 pct granular bainite, with a prior austenite grain size of approximately 50/xm (Figure 12). In the as-cooled untempered conditions, fine autotempered

Fig. 11--Optical micrographs of microstructures after slow-cooling at 8 ~ to simulate the 0.25T cooling rate in normalized 400 mm (16 inch) plate. (a) 2.25Cr-lMo steel, showing bainite and 40 pct poly- gonal ferrite, and (b) 3Cr-lMo-lNi steel, showing 100 pct bainite (etched in nital).

Fig. 12--Scanning electron micrograph of 3Cr-l.5Mo-0.5Ni Steel D in the slow-cooled and untempered condition, showing 100 pct granular ba- initic microstructure (etched in hot picric acid and nital).

M3C (cementite) precipitates (Figure 13) and films of re- tained austenite (Figure 14) are evident between bainitic laths. X-ray and magnetic saturation studies on Steel B have indicated that the proportion of such interlath austenite is approximately 17 pct in the as-cooled condition, compared to 5.5 to 7 pct in the as-quenched state. 29

On tempering at 700 ~ for 1 hour, fine accicular intra- lath carbides (--0.05 /xm in size) together with coarser lath boundary carbides (--0.25 /zm in size) were observed and

identified, using selected area diffraction, as cubic M23C6 (Figure 15). 32 No evidence of the hexagonal M2C or M7C3 or orthorhombic M3C carbides could be detected. This is in

stark contrast to 2 .25Cr- lMo steel where, at 700 ~ M3C precipitation can persist up to tempering times of 30 hours (Figure 16). 33 After 1000 hours at 700 ~ M23C6 precipi-

tates in the 3Cr-Mo-Ni steels are typically 0.5 to 1 bcm in size (Figure 17) with additional evidence of the Mo-rich M6C precipitates (Figure 18), and in the C and E steels VC

J. MATERIALS FOR ENERGY SYSTEMS VOL. 6, NO. 3, DECEMBER 1984 157

Fig. 13 Transmiss ion e lec t ron micrographs of s low-cooled 3Cr-1Mo-INi Steel B in the untempered condition, showing (a) bright field of ferrite (a) and_f'me autotempered cementite (0), (b) dark field of ce- mentite from (210)o reflection of (c) selected area diffraction (SAD) pattern, and (d) interpretation of (c). 32

Fig. 1 5 - - T r a n s m i s s i o n e lec t ron micrographs of s low-cooled 3Cr-l.5Mo-0.5Ni Steel D after tempering at 700 ~ for 1 h, showing (a) bright_field of ferrite and M23C~ precipitates, (b) dark field of carbide from (022) M23C6 reflection of (c) SAD pattern, and (d) interpretation of (c).32

Fig. 1 4 - - T r a n s m i s s i o n e lec t ron micrographs of s low-cooled (3Cr-l.5Mo-0.5Ni) Steel D in the untempered condition, showing (a) bright field of ferrite (a) and interlath retained austenite (3), (b) dark field from (200)v austenite reflection of (c) SAD pattern, and (d) inter- pretation of (c). 32

precipitates. The types of carbide in these structures were confirmed with extraction replica studies in the STEM using analyses developed by Titchmarsh 34 and Shaw. 35 Carbide types in Steels B, D, and E are shown in Table IV together with their respective compositions in atomic percent.

33 Fig. 16--Experimental data of Baker and Nutting showing the sequence of carbide formation during tempering of a normalized 2.25Cr-lMo steel.

DISCUSSION

On the basis of the TEM and STEM studies (Figures 15, 17, and 18), it is clear that the increased Cr and Ni content in the present 3Cr-Mo-Ni series of steels leads to a markedly ac- celerated tempering response compared to 2.25Cr-lMo steel. In particular, the iron and chromium-rich M23C6 car-

b ide replaces the less stable M2C, M3C and M7C3 carbides

158 VOL. 6, NO. 3, DECEMBER 1984 J. MATERIALS FOR ENERGY SYSTEMS

Fig. 1 7 - - T r a n s m i s s i o n e l e c t r o n m i c r o g r a p h s of s l ow-coo led 3Cr-lMo-lNi Steel B after tempering at 700 ~ for 1000 h, showing (a) bright field of ferrite and M23C6 precipitates, (b) SAD pattern, and (c) interpretation of (b). 32

Table IV. STEM Analysis of Carbide Precipitates in 3Cr-Mo-Ni Steels Tempered for 1000 hr at 700 oC

Composition (Atomic Pct) Type

Pct Cr Pct Mo Pct Fe Pct V Carbide

Steel B Fine ppts (0.1 /zm) 43.9 11.6 44.5 - - M23C6 Elongated ppts (1.0/zm) 24.7 32.7 42.6 - - M 6 C

Massive ppts (1.0/xm) 37.7 11.2 51.1 - - M23C6

Steel D Fine ppts 17.2 36.5 46.3 - - M6C Elongated ppts 36.3 17.5 46.2 - - M23C6 Massive ppts 46.0 7.3 46.7 - - M ; ~ 3 C 6

Steel E Fine ppts 8.8 4.0 1.2 86.0 VC Elongated ppts t0.6 45.8 43.6 - - M6C Massive ppts I 1.9 43.1 45.0 - - M6C

after only an hour of tempering at 700 ~ compared to approximately 400 hours in 2.25Cr-lMo steel (Figure 16). It is felt that specifically the rapid elimination of Fe3C (i.e., within 1 hour at 700 ~ compared to roughly 30 hours in 2.25Cr-1Mo) and the rapid precipitation of more stable alloy carbides on tempering is primarily responsible for the en- hanced hydrogen attack resistance in the 3Cr-Mo-Ni steels. This enhanced resistance to hydrogen damage can be seen clear ly in F igure 19 where the impact toughness 3Cr-lMo-lNi Steel B is compared to 2.25Cr-lMo Steel A in the prior hydrogen exposed and unexposed conditions

Fig. 1 8 - - T r a n s m i s s i o n e l e c t r o n m i c r o g r a p h o f s l o w - c o o l e d 3Cr-I.5Mo-0.5Ni-0.2V Steel E after tempering at 700 ~ for 1000 h, showing (a) bright field of ferrite, and M23C6 and M6C carbides, (b) dark field of (026) M6C reflection of (c) SAD pattern, and (d) interpretation of (c). 32

I ' I ' = -- 150~- _Stee_.~l B (._3.3Cr-!Mo-I Ni)

/ Quenched ond tempered / ot 650 =C/4 h lo Unexposed . . . . - -2 :

~j I /% H:~exposed , * ^ ~ - - i O0 - (6000c, o , o ~ . ~ - - ' e ' -

I �9 225Cr- IMo Steel

:~ / / ~ at 6500C/4h - ~. . s " / Z~I~" �9

C r ~ f H2 e~posed (550 C. 14 MPa)

200

150

I00 - , )

50

-200 -IO0 0 I00 200 Temperoture (=C)

Fig. 1 9 - - C o m p a r i s o n of Charpy V-no tch impact t oughnes s of 2.25Cr-lMo steel and experimental 3Cr-lMo-lNi Steel B, both quenched and tempered 4 h at 650 ~ prior to, and after, 1000 h high temperature exposure to gaseous hydrogen (550 to 600 ~ 14 to 17 MPa pressure).

(both steels were oil quenched and tempered 4 hours at 650 ~ Even though the 3Cr-lMo-lNi steel suffered a higher exposure (17 MPa hydrogen pressure at 600 ~ com- pared to 14 MPa at 550 ~ its toughness remained superior to the 2.25Cr-lMo steel, both in terms of a lower transition temperature and higher upper shelf energy.

The Cr + Ni combinations in the experimental steels also provide a potent effect in vastly increasing hardenability, as evident by the fully bainitic microstructures (Figure 11) following simulated normalizing of 400-mm thick plates (0.25T location). Mn contents, however, were deliberately limited to 0.5 pct as prior experience ~2 with a 3Cr-IMo-IN;

J. MATERIALS FOR ENERGY SYSTEMS VOL. 6, NO. 3, DECEMBER 1984 159

steel containing 1 pct Mn had indicated lower creep strength and problems with hydrogen attack susceptibility where, particularly after slow cooling, the transformation of high carbon retained austenite on tempering lead to a pro- longed stability of Fe3C carbides. Higher Mn contents are also known to promote banding, which again is detrimental to hydrogen attack resistance, and to increase susceptibility to temper embrittlement. 26,27,2s

Aside from generally improving temper embrittlement resistance, the Mo content of 1 to 1.5 pct provides for good elevated temperature strength and creep-rupture properties. This is further enhanced in Steels C and E by the precipi- tation of a vanadium carbide. Although in the past there has been some question about the creep resistance of 3Cr-Mo steels, 19'2~ the present series of alloys compared very favor- ably with 2.25Cr-lMo steel, at least over the testing range of temperatures and stresses (Figures 9 and 10).

Although not studied in the present investigation, the weldability of these steels should not cause problems with standard welding and pre-heat practices due to the relatively low carbon contents (i.e., O. 15 max wt pct). In addition, preliminary data on fatigue crack propagation behavior indi- cate a crack growth resistance similar to 2.25Cr- 1Mo steel.36

Thus, the present series of experimental 3Cr-Mo-Ni steels provides superior alternatives to 2.25Cr-1Mo steel for thick section pressure vessel applications involving hydrogen at elevated temperatures and pressures. Although evaluations have only been completed to date on relatively small heats (55 kg), the new steels appear to out perform 2.25Cr-lMo with respect to hardenability, strength, toughness and hy- drogen attack resistance and to be comparable with respect to ductility, temper embrittlement and creep resistance (Table V). There may be applications for lower tempera- ture (i .e. , below 500 ~ hydrogen service where the 3Cr-lMo-lNi Steel B may be preferred because of its very high toughness. 25 For service at higher temperatures, the

vanadium containing Steels C and E may be preferred be- cause of their superior creep strength.

Finally, the present series of steels compare very favor- ably with other recently developed 3Cr-Mo alloys for thick section, hydrogen service pressure vessel application, namely the 3Cr-I.5Mo-0.1V steels developed by Climax Molybdenum ~9'2~'23 and the 3Cr-lMo-0.25V, Ti, B steels developed by the Japan Steel Works. 24 As shown in Figures 20 to 22, all three series of steels out perform 2.25Cr-lMo steels. However, whilst the creep rupture (Figure 20) and temper embrittlement resistance of the three 3Cr-Mo steels is similar, the current 3Cr-Mo-Ni alloys appear to offer the best combinations of strength and toughness (Figures 21 and 22).

Fig. 20 - -Compar i son of the creep rupture strength of slow cooled and tempered experimental 3Cr-Mo-Ni alloys with 2.25Cr-lMo steel, 3~ the Climax Molybdenum 3Cr- l .5Mo steels, x9'21 and the Japan Steel Works 3Cr-lMo-0.25V, Ti, B steel. 24

Table V. Comparison of Mechanical Properties of New 3Cr-Mo-Ni Steels with 2.5Cr-lMo, Following Slow Cooling (8 ~ and Tempering at 700 oC (4 hr)*

Upper Creep Shelf Rupture Life

Yield Charpy 40 ft-lb at 200 MPa Code Alloy Stress UTS Elong.** RA Energy Transition (560 ~

MPa (ksi) MPa (ksi) Pct Pct J (ft-lb) ~ C (~ hr

A 2.25Cr- lMo 496 (72) 641 (93) 24 73 122 (90) - 13 ( 9) 457 B 3Cr- lMo- lNi 627 (91) 779 (113) 24 73 206 (152) - 1 0 5 ( -157 ) 269 C 3Cr-IMo-INi-0.2V 593 (86) 772 (112) 26 75 186 (137) - 25 ( - 13) 947 D 3Cr-l .5Mo-0.5Ni 655 (95) 799 (116) 22 73 203 (150) - 20 ( - 4) 212 E 3Cr-I .5Mo-0.5Ni-0.2V 676 (98) 821 (119) 21 74 176 (130) - 10 ( 14) 1253

*Tempering Parameter (K) = 20.04 **On 32 mm gauge length

160 VOL. 6, NO. 3, DECEMBER 1984 J. MATERIALS FOR ENERGY SYSTEMS

CONCLUSIONS

A new series of 3Cr-Mo-Ni steels has been developed to provide alternatives to 2.25Cr-1Mo steel for thick section pressure vessel applications involving high temperatures (up to 550 ~ and high hydrogen pressures (up to 18 MPa). The nominal compositions of these steels are 3Cr-1Mo-1Ni and 3Cr-l.5Mo-0.5Ni, with and without 0.2V. Based on evalu- ations on 55 kg heats the new steels are found to be superior to 2.25Cr-lMo steel with respect to hardenability, room temperature strength, impact toughness and resistance to hydrogen attack, and to be similar with respect to room and elevated temperature ductility, creep strength and rupture ductility and temper embrittlement resistance. In addition, their mechanical properties compare very favorably to the recently developed 3Cr-1.5Mo-0.1V Climax Molybdenum and 3Cr-lMo-0.25V, Ti, B Japan Steel Works steels, and actually are somewhat superior to these steels with respect to strength and toughness.

Fig. 21--Comparison of tensile strength of present experimental 3Cr-Mo-Ni alloys with 2.25Cr-lMo steel , 37'38 the Climax Molybdenum 3Cr-l.5Mo steels, ~9'2~ and the Japan Steel Works 3Cr-IMo-0.25V, Ti, B steel. 24

ACKNOWLEDGMENTS

This work was supported by the Department of Energy, Advanced Research and Development Fossil Energy Mate- rials Program, through the Oak Ridge National Laboratory (ORNL), under Subcontract 7843 with the University of California, Berkeley, under Union Carbide Contract W- 7605-eng-26 with the U.S. Department of Energy. The au- thors would like to thank R.W. Swindeman of ORNL for many helpful discussions, Dr. D. S. Sarma for transmission electron microscopy assistance, and T. George, R.K. An- ders, R. I. Huntley, and L. -H. Chan for experimental help.

Fig. 22--Comparison of the tensile strength/toughness characteristics of present experimental 3Cr-Mo-Ni alloys with 2.25Cr-lMo steel? 9 the Climax Molybdenum 3Cr-l.5Mo steels, ~9'2~ and the Japan Steel Works 3Cr-IMo-0.25V, Ti, B steel, z4 Note lower ductile/brittle transition tem- perature implies higher toughness.

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162 VOL. 6, NO. 3, DECEMBER 1984 J. MATERIALS FOR ENERGY SYSTEMS