reformer furnace 02

In service embrittlement of cast 20Cr32Ni1Nb components used

in steam reformer applications

D.M. Knowlesa, C.W. Thomasa,*, D.J. Keenb, Q.Z. Chenc

aMaterials Performance Technologies, P.O. Box 31310, Lower Hutt, New ZealandbPlant Reliability Solutions, P.O. Box 263, Carina, Qld 4152, Australia

cUniversity of Hong Kong, Pokfulam Road, Hong Kong, China

Abstract

Severe embrittlement has been experienced in a number of cast manifold components. This has manifested itself as cracking at tee to

manifold connections. Attempts to weld repair proved futile leading to concern about the integrity of the entire system. This experience

contrasts with similar components that have successfully remained in service for many years. The paper describes the investigations into

these failures and laboratory investigations into the properties of cast 20Cr32Ni1Nb alloys. Results indicate that variations in alloy chemistry

within the stated allowable range are sufficient to cause embrittlement.

q 2004 Elsevier Ltd. All rights reserved.

Keywords: Embrittlement; Creep–fatigue

1. Introduction

Steam reformer furnaces are at the front end of a number of

industrially important processes. These furnaces take a

supply feed of methane and steam and ‘reform’ them to

hydrogen and carbon monoxide which subsequently become

the basic building blocks in industries such as ammonia,

methanol, DR iron production and petroleum refining. A

typical reformer furnace consists of an array of vertical tubes

in a firebox. The smallest furnaces may have only ten such

tubes but the largest, with up to 700 tubes, are very significant

and capital intensive items of plant. These tubes contain a

catalyst and the feed gas flows internally from the top to the

bottom. Effectively, each tube behaves as a separate reactor.

At the bottom of the furnace, the various tubes are all

connected to a system of manifolds that collects the gas into a

single stream for distribution to further processing units. The

operating temperatures required in these furnaces are high.

Skin temperature of the reformer tubes isapproximately850–

950 8C and gas outlet temperatures are around 760–850 8C.

These temperatures and the need to operate reliably for

scheduled campaigns of possibly five years put huge demands

on the materials of construction used in these furnaces.

The present paper addresses problems encountered in the

outlet manifold system. Fig. 1 shows a schematic illustration

of a reformer and an outlet manifold system. Traditionally

manifold components have been manufactured from

wrought alloy 800H or 800HT. As systems have increased

in size, there has been a shift to the more economic and

nominally better performing cast 20Cr32Ni1Nb alloy which

has become an industry standard. This alloy is offered by a

number of manufacturers using various trade names but in

reality, there is little variation between them.

2. The alloy

The material is covered by ASTM Standard A351-94

where it is described as alloy CT15C. Reference to this

standard, however, is rarely made and the material is more

commonly identified by its various trade names such as

CR32W or KHR32C. This standard describes composition

and manufacturing requirements but makes no reference to

elevated temperature mechanical properties.

The material is essentially, a cast version of alloy 800.

Alloy 800 contains 20% chromium and 32% nickel with an

upper limit of 0.1% carbon. It is a solid solution alloy but

also contains small amounts of aluminium and titanium

which lead to the formation of carbides and sometimes,

a small amount of gamma prime g0 phase. The ‘H’ and ‘HT’

grades involve manipulation of grain size and minor

variations in the carbon, aluminium and titanium content.

0308-0161/$ - see front matter q 2004 Elsevier Ltd. All rights reserved.

doi:10.1016/j.ijpvp.2003.12.025

International Journal of Pressure Vessels and Piping 81 (2004) 499–506

www.elsevier.com/locate/ijpvp

* Corresponding author. Tel.: þ64-4-569-0027; fax: þ64-4-569-0431.

E-mail address: [email protected] (C.W. Thomas).

http://www.elsevier.com/locate/ijpvp

The adoption of a cast variant saw the use of similar levels

of the prime alloying elements, chromium and nickel.

Titanium and aluminium were not used in the cast variant.

However, in parallel with the development of spun cast

reformer tube alloys involving alloying with small amounts

of niobium, the manifold alloy similarly was alloyed with

approximately 1% of niobium to improve creep properties.

Table 1 lists chemical composition requirements of

ASTM A351 alloy CT15C and compares them with the

wrought alloy 800H. The similarities are clearly shown.

There are no standardised creep or stress rupture

properties for the cast 20Cr32Ni1Nb alloy. Instead,

designers are required to make use of data supplied by

manufacturers. Alloy 800 on the other hand, has been in use

for many years and stress rupture data is available from

a number of sources. Fig. 2 shows a comparison of

published stress rupture data for the cast material (based

on a manufacturer’s data) and the wrought alloy 800H [2]

(based on API 530 data). According to these data, the

expected life of the cast material is an order of magnitude

higher than that of the wrought material at typical design

stress levels around 10 MPa.

3. The problem

The problem that was encountered with this material was

severe in-service embrittlement. After only relatively short

periods in service, routine inspection at scheduled plant

outages led to the discovery of cracking at the weldments

Fig. 1. Schematic illustration of a reformer furnace [1].

Table 1

Chemical compositions of CT15C and alloy 800H

%C %Cr %Ni %Si %Mn %Nb %Ti %Al

CT15C 0.05–0.15 19–21 31–34 0.5–1.5 0.15–1.5 0.5–1.5

800H 0.1 max 19–23 30–35 1.0 max 1.5 max 0.15–0.6 0.15–0.6

Sample 0.11 19.74 31.37 1.03 0.90 0.98

Fig. 2. Comparison of stress rupture data for cast alloy CT15C and wrought alloy 800H.

D.M. Knowles et al. / International Journal of Pressure Vessels and Piping 81 (2004) 499–506500

connecting the main manifold arms to the bull-T. The bull-T

typically connects the arms of the manifold to the main

transfer line. Typical schematic manifold/bull-T arrange-

ments are shown in Fig. 3. Fig. 4 shows a bull-T in position

below the furnace. This problem was encountered at two

independent plants.

In service, the manifold undergoes significant thermal

expansion. The tube to manifold pigtail connections,

despite being deliberately flexible to account for thermal

expansion loads are numerous and, in acting together, are

capable of exerting significant system loads on the

manifold. These thermal loads lead to significant bending

applied to the manifold arm to bull-T weldment and it is

not uncommon for cracking to be found at this location.

The most common failure mode in these manifold

systems is in fact, creep–fatigue at the welds and this

was found to have occurred at two independent plants.

Fig. 5 shows the sort of cracking that was observed.

Based on the stress rupture data shown in Fig. 2, an

effective life of 5 years and a service temperature of

760 8C, an approximate mean stress of 45–50 MPa is

implied. This is significantly higher than pressure based

hoop stresses which are typically below 10 MPa and upon

which design is based and indicates the relatively high

level of thermal stress.

What was not expected, however, was the extreme

brittleness of the parent bull-T material. Attempts to grind

out and re-weld the damage simply led to the generation of

more cracks as can be seen in Fig. 6.

It is not uncommon for in-situ solution anneal heat

treatments to be conducted on these materials to improve

weldability. The temperature required, however, is in excess

of 1100 8C and a recent recommendation was that this

should be increased to 1200 8C. Such heat treatments can be

readily done for small areas in the immediate vicinity of the

weldment. This did not, however, provide confidence that

the remaining parts of the manifold system were sufficiently

ductile to ensure their integrity when returned to service.

Fig. 3. Schematic illustration of typical manifold and bull-T arrangements.

Fig. 4. Typical bull-T. Fig. 5. Creep fatigue crack.

D.M. Knowles et al. / International Journal of Pressure Vessels and Piping 81 (2004) 499–506 501

This, led to the decision in both cases investigated, that

temporary repairs should be made until such time that

replacement components could be installed. One of the bull-

Ts was made available for metallurgical investigation. It had

been in service for approximately 5 years at service

temperatures of approximately 760 8C. The analysis of

this material is included in Table 1.

4. Mechanical testing

The issue of concern was apparent extremely low

ductility that led to an inability to weld the bull-T without

cracking. Consequently, the mechanical testing undertaken

to date has concentrated on toughness using Charpy impact

testing on standard 10 £ 10 mm samples. In addition, a

series of tensile tests have also been undertaken. The testing

has been conducted on the material removed from service

(as received) and after a solution annealing heat treatment.

For the as-received material, a series of Charpy impact tests

were undertaken at a series of temperatures approaching

the service temperature. The intention was to determine if

the apparent loss of ductility was a low temperature

phenomenon or if toughness was also compromised at

service temperatures. Tensile testing similarly was under-

taken at room temperature and at 800 8C to simulate service

conditions.

The solution annealing heat treatment involved holding

at 1100 8C for 3 h followed by air cooling. The small

laboratory specimens were removed from the furnace and

cooled relatively rapidly in air.

The Charpy impact test results are listed in Table 2 and

illustrated in Fig. 7. The tensile test results are contained in

Table 3.

The Charpy data exhibited some scatter. However, while

there was a modest increase in toughness with increasing

temperature, the toughness as revealed by impact testing

remained low even at temperatures close to the operating

temperature. The tensile data also revealed extreme

brittleness at room temperature for the as-received ex-

service material with no measurable elongation on the test

piece itself. The stress–strain curve for this sample showed

a plastic strain of less than 0.4%. Only one test was

undertaken for this condition because the material was so

brittle, the duplicate specimen failed during machining.

The as-received material however, had significant ductility

at service temperatures. In fact, the elongation and reduction

Fig. 6. Cracking after and during attempted weld repair.

Table 2

Charpy Impact Results

Condition Temperature (8C) Impact energy J

Annealed 23 32

Annealed 23 38

Annealed 23 26

Ex service 23 6

Ex service 23 9

Ex service 23 7

Ex service 105 14

Ex service 200 11

Ex service 295 13

Ex service 320 20

Ex service 390 10

Ex service 590 12

Ex service 710 25

Ex service 760 10

Fig. 7. Charpy impact data.

Table 3

Tensile test results

Condition Temperature

(8C)

UTS

(MPa)

0.2%Proof

(MPa)

% Elong % ROA

Sol annealed 20 340 158 15 13

Sol annealed 20 16 15

Sol annealed 800 180 91 39 45

Sol annealed 800 171 80 39 54

Ex service 20 200 177 0 0

Ex service 800 171 80 51 55

Ex service 800 174 82 34 33


of area results were as high or higher for the as-

received material than the solution annealed test pieces at

800 8C.

5. Metallography

The microstructure of 20Cr32Ni1Nb consists of an

interdendritic network of primary carbides in an austenitic

matrix. The microstructure is illustrated in Fig. 8. Despite

the interdendritic eutectic carbides and the relatively low

carbon content of this material, the austenite matrix is

typically supersaturated with carbon and some fine intra-

dendritic carbides can be seen in the matrix.

The density of intra-dendritic carbides typically

increases dramatically when the material first enters service.

Fig. 9 shows a similar sample to that shown in Fig. 8 except

that it has been held at 800 8C (approximate service

temperature) for 24 h.

The number and density of intra-dendritic secondary

carbides is extremely high. It is these carbides and the

complexity of the interdendritic eutectic carbides that have

been attributed with generating the good creep properties of

similar alloys [3].

In the ex-service material, the relatively long term aging at

approximately 800 8C had led to the agglomeration and

dissolution of many of the intra-dendritic carbides. A

significant number however, remain as can be seen in Fig. 10.

Fig. 9. 20Cr32Ni1Nb alloy after 1 day at service temperatures.

Fig. 8. As-cast 20Cr32Ni1Nb.

Fig. 10. 20Cr32Ni1Nb alloy ex-service.


The in-serviceexposure has not led towholesale changes in the

microstructure. It is considered significant, however, that the

network of primary carbides appears continuous in many areas

and has adopted a two phase appearance.

After solution annealing (3 h at 1100 8C), the microstruc-

ture has again not changed in any major way. The density of

intra-granular secondary carbides remained similar to the ex-

service material except a precipitate free zone has developed

adjacent to the primary carbides. In addition, the dual phase

nature of the eutectic carbides was removed and the strings of

interdendritic precipitate that were present in the ex-service

material tended to break up into discreet particles (Fig. 11).

As part of an on-going program to understand the

microstructure of this alloy, the ex-service material has

been examined using a scanning electron microscope.

The dual phase nature of the primary carbides was

evident as illustrated in Fig. 12 which shows an example

of primary eutectic precipitation viewed using the back

scatter detector to highlight atomic weight differences.

Analysis of these precipitates using an energy dispersive

X-ray analysis system in conjunction with the SEM

showed the two phases within the precipitates to be

strongly segregated (Fig. 13). One was chromium rich

while the other was essentially free of chromium and

contained niobium and silicon. This silicon/niobium rich

phase was not found in the as-cast or solution annealed

materials.

6. Discussion

The Charpy impact testing confirmed the brittle nature

of the ex-service material. There was a modest improve-

ment in toughness with increasing temperature but the

impact energy of the ex-service material was always

below that of the solution annealed material. The low

ambient temperature tensile tests also showed the ex-

service material to be extremely brittle with no measur-

able elongation being recorded on the test pieces. It was

of interest however, that in tensile tests at 800 8C, tensile

ductility of ex-service material was high. The solution

annealing heat treatment re-established a significant level

of toughness. These mechanical test results are consistent

with the observation that the original bull-T was

extremely difficult to weld. However, the relatively high

tensile ductility at operating temperatures suggest that

integrity at operating temperature is not an issue. Thermal

loads at start up or shut down however may lead to

cracking while the bull-T is relatively cold.

The metallographic examination revealed no immedi-

ately obvious reason for the embrittlement. More in depth

examination however, revealed the presence of a silicon and

niobium rich phase in the interdendritic precipitates. These

materials are part of an on-going investigation to charac-

terise the microstructure and establish the influence of these

phases on material properties. Similar phases in

20Cr32Ni1Nb have however been identified by other

workers[4,5] who have identified the silicon rich phase as

‘G-phase’ reported to be Ni16Nb6Si7. In similar but higher

carbon alloys such as HP50Nb reformer tube materials [6],

the silicon rich phase was identified as a silicide having an

h-carbide (M6C) structure.. In all these cases, extreme

brittleness at ambient temperatures have resulted.

It is concluded therefore that the problem of brittleness in

the outlet manifold components examined was caused by

Fig. 11. 20Cr32Ni1Nb alloy ex-service and solution annealed.

Fig. 12. SEM backscattered electron image of primary eutectic interden-

dritic precipitates in ex-service material.


the formation of intermetallic niobium rich silicide phases.

The observation that improvement in weldability can be

obtained by solution annealing is consistent with the

observation that these silicide phases were not present in

the solution annealed samples.

Based on the observations made, the formation of these

deleterious phases can be considered a normal consequence

of in-service aging. This is clearly an undesirable situation.

It has been suggested [4] that the formation of the silicides

can be controlled if not prevented, by ensuring that the

niobium level is maintained below that necessary to

stoichiometrically accommodate all the carbon as NbC,

i.e. the wt% ratio of niobium to carbon should be held below

7.7. The observation of silicides in HP50Nb materials where

the carbon content is higher at 0.4–0.5% would suggest that

this is not sufficient to prevent the formation of silicides.

It would appear that niobium is a prerequisite for the

formation of silicides. It is therefore questioned why

niobium needs to be used in this application at all. Alloying

of cast alternatives to alloy 800 with niobium appears to

have been adopted by manufacturers because of experience

with tube alloys and an apparent improvement in creep

properties that this produces. Experience with HP50 alloys

[7] has shown that in aged materials, the stress rupture

strength of tube materials are not significantly better than in

the niobium free versions, i.e. any strength advantage

gained by alloying with niobium is soon lost once the tubes

have entered service, probably because of the formation of

intermetallic silicides. In addition, in manifold components,

wall thickness is not important and strength can therefore be

obtained by design rather than material strength.

However, more work is clearly required before the use of

niobium alloyed materials should be condemned in this

application. For example, a benefit that has been observed in

the performance of aged niobium alloyed materials is their

relatively good creep ductility [7]. This property is

important in maximising resistance to creep fatigue, which

after all, was the problem that initiated this investigation in

the first place. The good high temperature ductility of the

20Cr32Ni1Nb material is suggested by the high tensile

ductility of the ex-service material which was superior to the

solution annealed material.

7. Conclusions

20Cr32Ni1Nb cast material has become an industry

standard for reformer furnace outlet manifold components

In-service aging of this material, however, leads to serious

embrittlement that yields the material unweldable. This

problem can be at least mitigated by solution annealing at

temperatures in excess of 1100 8C, but this is not practical

for complete outlet manifold systems which remain

embrittled. Furthermore, solution annealed materials are

likely to re-embrittle on further exposure.

The cause of the embrittlement is the formation of

niobium rich silicide intermetallics which have also been

observed in closely related HP50Nb reformer tube alloys.

The presence of niobium appears to be a prerequisite for the

formation of these intermetallics. It is therefore suggested

that niobium free grades are developed to avoid the problem

of in service embrittlement.

Fig. 13. Element map showing distribution of silicon, chromium and niobium in interdendritic precipitates in ex-service material.


References

[1] Rostrup-Nielsen JR. Catalytic Steam Reforming. Berlin: Springer;

1984.

[2] API 530 Calculation of heater tube thickness in petroleum refineries,

American Petroleum Institute, 1220 L Street, Northwest Washington,

DC 20005.

[3] Hou W-T. Mater Sci Technol 1985;1(5):385–7.

[4] Hoffman JJ, Gapinsky GE. Ammonia Plant Safety. AIChE 2002;42:

10–21.

[5] Shibaski T, Mohri T, Takemura K. Ammonia Plant Safety, AIChE

1994;34:166–76.

[6] Thomas CW, Stevens KJ, Ryan MJ. Mater Sci Technol 1996;12:

469–75.

[7] Thomas CW, Tack AJ. Proc Int Symp. Case Histories on Integrity and

Failures in Industry, Milan Sep 1999;28:1.


reformer furnace 02

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