EXTREME C O N D I T I O N PROCESSING
Effects of Hydrogen on Properties of Metals D. D. PERLMUTTER'
AND
B. F. DODGE
Yale University, New Haven, Conn.
W
ITH the exception of the published preliminary results of
this investigation (60) and two German articles by Naumann (26, 27), no systematic measurements of the effects of hydrogen on metals are available covering the range above several hundred atmospheres pressure. As industrial needs begin to approach higher pressures, the necessity for design data becomes more than just an academic question, providing immediate applicability for what would in any case be of great theoretical interest. Specimens were exposed t o hydrogen gas atmospheres at pressures between 7500 and 60,000 pounds per square inch a t room temperature and to pressures of 15,000 and 30,000 pounds per square inch temperatures up t o 500" C. The time of exposure was varied considerably between 6 hours and 60 days so that information would be available as t o the rate of the embrittlement process. Only six alloys of the more than fifty tested were especially prepared for this work; the rest are all commercially available materials. Theoretical Considerations
A discussion of the interaction between a gas and a liquid, difficult though i t may be, is greatly simplified by the relative ease with which a liquid is characterized; but this is not the case in considering a solid metal. T o indicate only part of the complexity of such a problem, recall that commercial grade metals are never pure in a chemical sense. Furthermore, they are usually composed of more than one phase, seldom without inclusions themselves. The lattice of the purest metal or alloy is interrupted by atomic imperfections such as dislocations, vacancies, and mosaic-block boundaries, and not too infrequently by somewhat larger but still tiny cracks or flaws. Even if we avoid compounding the complexity by superimposing the still largely unknown effect of surface condition, it IS not difficult to understand why investigations in this field have been a t times confused. The fundamental work in this field is primarily concerned with adsorption, solubility, diffusion, and permeation. The subject of hydrogen solubility in metals has been aided greatly by the research of Sieverts and his coworkers (documented in do), mho have investigated a large number of pure metals over wide temperature ranges; nevertheless, the data are of very limited use for our purposes for two reasons. The first is due to the considerations of metallic structure just outlined. I n order that solubility data be reproducible a reference condition is essential; the one most conveniently reproduced is the fully annealed state. Unfortunately for engineering use most metals are heat treated or worked out of an annealed state for other desirable properties. The second major limitation is pressure range: Sievert's data do not go above l1/* atmospheres pressure. The relation called Sievert's law, the proportionality of solubility t o the square root of pressure, is not exact for all metals even over the range of 11/2 atmospheres where it was formulated, a fact which make8 extrapolation to high pressures a very gross approximation. In contrast to solubility, data on diffusion and permeation are less organized and consistent but closer t o the range of variables of our interest (9). Until the importance of surface reactions 1
Present address, Esso Research and Engineering Co., Linden, N. J.
May 1956
was understood, many investjgators confused their measurements and reported as diffusion in the metal what they actually measured as the over-all permeation process which includes the surface reactions at entrance and exit from the metal. Diffusion Relations
A slab-shaped test specimen completely immersed in a gas atmosphere provides a special geometric case easily handled mathematically. If a slab of thickness, h, is surrounded by diffusing gas of constant concentration, the differential equation t h a t relates concentration in the metal, C, to time, 1, and distance through the metal, x, is aC/& = D(d2C/dx2) (1) where D is the diffusivity of the gas in the solid (assumed to be constant). With initial concentration of gas within the metal equal to Cs, and the equilibrium concentration, CO,the boundary conditions are
ci co C(t,h) = co
C(0,z) =
(2)
C(t,O) =
(3)
(4)
and the solution for fractional saturation, f,is
c - ci f= ms=
When the fractional point saturation is integrated with respect to volume, this gives fractional average saturation, F , dependent only on time, in terms of 'E, the average concentration:
There are two special case limits for this equation especially worth mention. As t becomes small (at the start of a run) the equation simplifies to
(7) For very large values of t
F
= 1
8 - 2-?r
exp
Both these simplifications are important for interpretation of the results presented here. Surface Effects
However, a knowledge of surface phenomena is a t least as important as an understandjng of effects internal t o the metallic
INDUSTRIAL AND ENGINEERING CHEMISTRY
885
ENGINEERING. DESIGN. AND PROCESS DEVELOPMENT ~~~
structure. Before all else the hydrogen molecule meeting the metal, if it is to do more than merely bounce off, must interact with the surface, The over-all picture of surface catalysis and gas adsorption on metal surfaces has undergone considerable alteration in recent years. Current theories, not yet fully developed, involve: magnetism (35); resonating d-shell electrons (12, 31) ; electron transfer, molecular dipoles, and work function (6, 36); and crystal structure and steric orientation ( 4 3 ) . This richness in h - potheses should not be interpreted as utter confusion; on the contraiy, it demonstrates the degree of growth in this very fertile field. There is experimental as well as theoretical backing for each approach, and it is likely that when a more comprehensive treatment of the whole field is achieved. it will include elements of all these contributions. Since 1935 (33)it has become increasingly evident that what one ordinarily considers a clean metal surface is in fact covered by adsorbed atoms or molecules and, \That is even more important, that it is precisely these contamination layers that offer the first obstacle to a colliding atom. Hence, except where special care is taken, the “ordinary” case of gas-metal interaction such as hydrogen absorption is in reality a complex adsorption on a t least an oxide film formed by air oxidation and often even more complicated. The character of this all-important film depends on many factors-e.g., temperature and surface roughness-and its porosity will determine hos- the metal mill respond to high pressure hydrogen gas (16, 44).
Embrittlement and Attack Of the two classifications, embrittlement and attack, the former is the more fundamental, involving interactions of a subtler kind than the chemical changes that constitute attack. While attack can be explained easily as the radical property changes accompanying chemical reaction, the mechanism of embrittlement is still largely speculative; current ideas dram- primarily on what is known of dislocation theory aiid interstitial solution. Except for a few isolated measurements (7, S d ) , embrittlement caused by high pressure hydrogen gas was a virtually unexplored field a t the time this investigation was begun by Van S e s s and Dodge, but literature references on acid pickling embrittlement and electrolytic embrittlement were plentiful. Indeed, the literature of hydrogen embrittlement is so voluminous that it is only possible t o give a few references here (3, 5 , I S , 25) and refer the reader to further bibliography sources: U. S. Department of Commerce, circular 90.511 of the Sational Bureau of Standards Kith 1191 references, and Smith’s textbook (40). In 1952 a report appeared of extensive investigations carried out in an industrial laboratory (8),but no detailed results \\-ere published. JT7hat information JYas released was confirmation of the early reports of Bridgman (7) and OF Poulter and Uffelman (38) that in very short periods of time, sometimes minutes, pressure vessels satisfactory for containing 7000 atmosphere oil pressure, failed in hydrogen service under only 2000 atmospheres pressure. On the other hand, the phenomenon of chemical attack by hydrogen gas has been well explored. although in a limited pressure region: Max. Temp.,
c.
cox ( l o p Inglis and Xndrews ( 1 7 ) Sarjant and hIidclleham ( 3 4 ) Jacque (18) Naumann 126. $?) Nelson (28; 29)’ . Van P\-ess and Dodge (60)
700 500 750 550 600 600 500
Max. Press., Atm. 200 250 250
5 Primarily reviem papers T%ith some plant data reported. entries the maxima refer t o t h e plant d a t a .
886
1.50 970 700 2000 F o r these
In addition, much has been published on the effects of hydrogen-nitrogen ammonia synthesis mixture: Max. Temp., O
Wheeler (61 ) a Vanick (45-49) Kosting ( 2 2 ) Maxwell ( $ 4 ) Ihrig ( 1 6 ) Schuyten (35)b
c.
Max. Press., Atm.
750 500 300 500 590 600
1
100 600 1000 1000 130
a D a t a listed are for Wheeler’s experimental work i n which ammonia was decomposed in situ to produce the mixture. I n t h e ~ a m earticle Wheeler reports some plant results a t 50Oo C . and 100 a t m . b Primarlly revieiv papers with some plant d a t a reported. For these entries the maxims refer to t h e plant d a t a .
Theories of Embrittlement The task of relating fundamental developments in solid-state theory t o experimental embrittlement findings measured by overall mechanical properties is not a simple one; yet, because such connections are of prime importance, several attempts to unify all experimental data have been made. One long series of such developments has been built around the belief that larger-than-atomic imperfections are responsible. Kazinczy’s theory of Griffin In essence, Smith’s “rifts” (&I), cracks as initiating agents ( d l ) , and Zapffe’s “planer p r e s ~ u r e ’ ~ theory (62) all depend on voids of microscopic dimensions. In these holes. it is argued, hydrogen atoms diffusing through the metal can collect, recombine to diatomic molecules, and exert very great disrupting gas pressures. The second school of thought on the subject relies on the relatively recent findings of dislocations and vacancies in atomic lattices. This is more than merely a difference of degree, however, for the mechanism of proton concentration and metal reactions on an atomic scale can be of a very different kind, involving atomic diffusionlilre jumps and electron transfer considerations (38). The major role played by dislocations in regulating slip and the obvious association of slip and ductility have provided the links for this explanation. The dislocation concept bears a primary relation to slip and ductility behavior; in fact, it was to explain plastic effects and their numerical load requirements that the new slip mechanism m-as first postulated by Taylor (43) and Orowan (SO), Essentially, the view proposed (and now largely accepted) was that macroscopic slip is actually dislocation movements on an atomic scale. Cottrell has pointed out that dislocations would attract foreign atoms whenever their concentration at the dislocation site LYould relieve the local stresses. This is especially true of the more easily mobile interstitial atoms such as carbon, nitrogen, or hydrogen. It is reasoned, by analogy to carbon solution in iron. that protons strategically positioned a t dislocation sites could interfere with metal slippage and cause embrittlement (4, 15).
Methods of Ductility Measurement The qualitative concept of ductility as the ability to deform without rupture is hopelessly incomplete for numerical measurement unless it is coupled with a reference test. That is, if results are t o be meaningful, one must know by what ductility index they were measured. Among the various tests used for hydrogen embrittlement, one may cite, for example, the standard tensile test (37, 4 1 ) , Erickson cup test ( 2 3 ) , alternating stresa bend test (go), and impact test (42). test (E?), Each of these has been used with some measure of success, but it is important to note that there is no a priori reason to except agreement between any two; also, none of these (or in fact, any idealized laboratory test) can be used directly to predict service conditions (14). K h a t then can be expected of such experiments? T o be of use for engineering design, laboratory
INDUSTRIAL AND ENGINEERING CHEMISTRY
Vol. 48, No. 5
EXTREME CONDITION PROCESSING tests must at least be extensive enough t o indicate the relation between the particular test results used and some of the physicalchemical variables of the system considered. Under such circumstances, a scale of relative detrimental effects can be built for design purposes, and a link can be established between mechanical measurement and metallurgical theory.
Cupalloy, Westinghouse: 0.7 Cr, 0.10 Ag, 0.015 max. Pb, 0.04 max. P, 0.15 max. other, balance Cu. H a nes Alloy 25: 0.15 max. C, 20 Cr, 10 Ni, 1.5 max. Mn, 1 max. 15 eba.ls.nce . cO. - ~ ~,., - - W. R Hyperco, Westinghouse: 35 Co, 1.5 other, balance Fe. Hypernik, Westinghouse: 50 Ni, 0.90 Mn, 0.50 Co, 0.45 Si, 0.60 other, balance Fe. Incoloy, International Nickel: 0.05 C, 21.2 Cr, 34.8 Xi, 0.42 Cu, 0.71 Mn, 42.5 Fe, 0.43 Si, 0.008 S. Low carbon steel, U. S. Steel: 0.016 C, 0.011 Cu, 0.02 Mn, 0.007 Si, 0.005 P, 0.021 S. Maxel shank steel, Crucible: 0.50 C, 0.65 Cr, 1.25 Mn, 0.18 Mo, 0.08 S. Nichrome: 0.06 C, 16 Cr, 60 Ni, 24 Fe. Nichrome V: 0.06 C, 20 Cr, 80 Ni. Nickel silver: 10 Ni, 65 Cu, 25 Zn. Nifer, Metals and Controls Cor : low carbon steel, clad on each side by grade 330 nickel. &adding thickness, 10% on each side. Puron, Westinghouse: 0.005 C, 0.0015 Xi, 0.001 Co, 0 . 0 4 N, 0.04 0, 0.001 P, 0.001 Si, 0.003 S, 99.95 Fe. Rex M-2 tool steel: 0.83 C, 4.15 Cr, 1.9 V, 6.4 W, 5 Mo. Titanium-manganese alloy, Rem-Cru: 0.1 max. C, 7 . 3 Mn, 92.6 Ti, 0.03 N. Vitrenamel, U. S. Steel: 0.03 C, 0.07 Mn, 0.005 P, 0.031 S, 0.006 Si. Special alloy No. 1: 10 Cr. 90 Fe. No. 2: 15 Cr; 85 Fe. No. 3: 20 Cr, 80 Fe. No. 4: 10 Cr, 90 Fe, 0.15 C. No. 5: 15 Cr, 85 Fe, 0.15 C. No. 6: 20 Cr, 80 Fe, 0.15 C.
&. --I
I
.
d Figure 1.
Pictorial diagram of bend tester
Specimen held in tester Set screw for holding rpedmen in place C. Metal stop positioned for 90' sweep of specimen D. Arc showing total 180' sweep of specimen E. Arm used to bend tester and specimen A.
B.
For room temperature embrittlement tests, a repetitive bend test on a strip of metal sheet proved t o meet these requirements, while hollow-drilled standard tensile specimens could be used t o simulate high temperature-high pressure service conditions and give indication of chemical attack when it mas present, as well as hydrogen embrittlement. The details of equipment design were described and illustrated in the earlier report on apparatus, procedures, and preliminary results (50).
Bend Tests-Ro
o m Temperature Exposure
Any mechanical test is subject t o a combination of errors, ranging from limitation on the sensitivity of the measuring devices to the statistical variation in sampling metallic properties. The ductility recorded in this test is the number of 180" bends inch. required t o rupture a sample of sheet metal 4 X '/z X To make i t possible to reject atypical results several tests (usually four) were made and their results averaged. I n addition, a probable deviation mas calculated for each average t o facilitate comparison with t h e control specimens. The bend tester used is illustrated in Figure 1 and some typical results are shown in Table I. For convenience of presentation, the results are grouped according to the chemical classification of the metal. This should not be interpreted as necessarily indicating uniform behavior within a group, for as pointed out earlier, chemical composition is only one of many variables t h a t affect a metal's resistance t o high pressure hydrogen gas. Chemical composition data (%) on the less common alloys mentioned in this section and the next are as follows: Alfer, Metals and Controls Corp.: low carbon steel clad on each side by aluminium containing 1.2 Si; cladding thickness, 10% on each side. Brmco iron: 0.03 C, 0.08 Cu, 0.02 Mn, 0.009 P, 0.026 S. Cadmium bronze: 99.1 Cu, 0.9 Cd. Carilloy T-1, U. S. Steel Co.: 0.14 C, 0.60 Cr, 0.91 Ni, 0.33 Cu, 0.83 Mn, 0.23 Si, 0.40 Mo, 0.05 V, 0.0055 B, Chromax, Driver-Harris: 0.06 C, 20 Cr, 35 Ni, 45 Fe.
May 1956
Alloys omitted from list are common enough to be found in any standard reference book.
Table 1.
Typical Bend-Test Results
Hydrogen Gas Exposure Conditions Lb./sq. in. Days
Av. No. Bends
Probable Deviation
No. Specimens
Low Precision Results : Cold-Rolled Armco Iron (0.015% C, Mn, 0.001% Si, 0.00370 P,0.02% S) Control, as received 7,300 7,800 15,100 29,800 30,700 52,000 52,000 52,000 52,000
40 19.9 9.8 2 2 2 2 4
10
62.0 20.2 57.0 63.9 26.8 62.8 59.0 63.5 27.5 15.8
1.5 3.7 3.3 1.3 7.8 3.1 7.4 4.1 8.8 3.0
0.02% 12 4 4 4 5 9 5 3 6 6
Av. Precision Results: 35% Cold-Reduced Haynes Alloy 25 (0.15% Max. C, 1.5% Max. Mn, 1% Max. Si, 10% Ni, 20% Cr, 15% W, 3% Fe, Balance Co) Control, a s received
26.2
15,100 15,400 30,700 52,000 52,000 52,000 52,000
17.0
9.8 4 2 2 4
10 18.9
16.5 17.4 11.0 8.0 3.0 0.2
2.2 2.4 4.5 4.4 2.4 1.8 0.48 0
High Precision Results : Cold-Rolled Nichrorne V (0.06% C, 80% Ni, 20% Cr) Control, as received 18.0 0 3 29,800 51,800 52,000 52,000 52,000
2 60 2 4
13.0 2.5 11.0
10
6.0
INDUSTRIAL A N D ENGINEERING CHEMISTRY
10.0
0
..
0.30 0.05 0
3
1 3 2 2
887
ENGINEERING, DESIGN, AND PROCESS DEVELOPMENT Group 1 : Nickel and High Nickel Alloys. Nichrome V Chromax hlonel Hypernik
Kickel Incoloy Inconel Nichrome’
With the exception of Incolog and Chromax which were used in the early shorter-time runs and Hypernik, the metals of this group were all embrittled. However, for Inconel and Xichrome exposures up to 60 days a t the very high pressure of 52,000 pounds per square inch were required to clarify the question.’ The
of cold work w i s very evident here. Khile the annealed material shovied considerable resistance and was only moderately embrittled by 52,000 pounds per square inch gas after 60 days, the 35% cold-reduced specimens lost all their ductile character in 20 days and were markedly changed even after only 2 days. The third condition of cold work, 50% cold reduced, was already rather brittle in the as-received condition. The tests with this material served therefore only as a qualitative check. The same was true for molybdenum. The metal samples received were too brittle t o be used for extensive work and provided merely a yes-or-no result. On the basis of one run at 52,000 pounds per square inch for 60 days, molybdenum is embrittled. Group 4 : Stainless Steels and High Chromium Alloys. Stainless Steel Types 410, 420, 430, 321, 304 Iron-Chromium Alloys: 10, 15, and 20% Cr, with and without 0.15% c.
I
‘0
4
9
16
TIME IN DAYS -SQUARE
Figure 2.
25
36
49
64
ROOT SCALE
Ductility of Monel e x p o s e d to hydrogen g a s at room temperature
results found for Hypernik are believed to be due to a protective surface. The two most susceptible metals were Monel and Nichrome V-both were promptly embrittled even a t more moderate pressures. Because the hydrogen effect on blonel is an excellent example for further discussion, the data are presented here in graphical foim (Figure 2). The abscissa scale chosen has the dual advantage of conveniently compressing the time scale and of giving a straight line a t low values of time when the reaction is diffusion controlling, as predicted by the mathematics of Equation 7. Group 2 : Copper, Zinc and Their Alloys. Nickel Silver Yellow Brass Muntz Metal Cadmium Bronze
Copper Zinc Cupalloy
Viewed as a group the copper-zinc alloys arc quite resistant to hydrogen gas embrittlement. Only copper and cadmium bronze (lye Cd) showed slight changes after the severest exposure conditions. The others were wholly unaffected, as far as our measurements are concerned. Group 3 : Titanium, Cobalt, Aluminum, Molybdenum, and Their Alloys. Titanium Haynes Alloy 25
Aluminum Molybdenum
The results with three grades of commercial titanium and one titanium-7% manganese alloy were among the most optimistic. Over a n-ide range of severe exposure conditions-up t o 60 days a t 52,000 pounds per square inch-no statistically significant embrittlement was found. Several tests run on commercial grade 2-5 aluminum n w e also negative, bearing heavily on the interpretation of the results with aluminum-clad steel to be discussed later. The cobalt-base material used was Haynes alloy 25, which was obtained and tested in each of the following conditions: annealed, 35% cold reduced, and 50% cold reduced. The effect
888
By its presence alone, chromium does not provide protection against embrittlement by hydrogen. HotT-ever, when it is used t o create a stable chromium-oxide surface film as in stainless steel Types 430 and 321, it can definitely impart resistance to an alloy depending on service pressure and intended time of exposure. As an example consider Figure 3 where ductility versus the square root of time is plotted for stainless steel type 430. The induction period, presumably the time needed for the hydrogen t o penetrate the surface oxide film, grows as the pressure decreases until a t pressures below 7500 pounds per square inrh it exceeds 40 days and becomes impractical to measme. A t very high pressures of about 60,000 pounds per squaie inch, this time of piotection is only 1 day. A series of especially prepared Fe-Cr alloys provided additional evidence for this picture. Their chromium content, 10 to 20%, was sufficient t o provide a glazed finish which showed the familiar induction period. They were also similar to the stainless steels in respect to severity of embrittlenient. It must be emphasized here that until now we have been discussing only room temperature embrittlement. The effect of chromium addition on high temperature attack (decarburization) is taken up later. The 13% chromium stainless used (Type 420) did not have sufficient chromium to form as stable a chromium-oxide surface film and, as a consequence (Figure 4), embrittlement began to take place immediately upon exposuie The influence of pressure is to speed up the rate of embrittlement considerably, but perhaps the most significant featuie of these data is the limiting embrittlement value. Although it was quickly affected, Type 420 stainless never lost its ductile character completely as did other metals. Among the stainless steels, Type 321 also shored great residual ductility. The embiittlement limit is probably associated with a saturated state with respect to hydrogen. Group 5 : Iron and Steel, Ferrous Alloys, and Clad Steels.
Low Carbon Steel Carilloy T-1 Rex M-2 Titanium Steel 3laxel Shank Steel Puron Vitrenamel
AISI 1020 AISI 1010 AIS1 1095 Nifer Alfer Galvanized Iron Armco Iron
This last group of metals, probably the most important, is also the most complex to systematize. Alloying element contents cannot be of sole importance, for such chemically similar metals for example as Armco iron and low carbon steel (both commercial grades of pure iron) gave radically different results. Although low carbon steel could not be embrittled under extreme pressure for long periods, Armco iron samples were erratically but rapidly affected most severely. T‘itrenamel and Puron, the other relatively pure irons tested, were like rlrmco iron in their behavior. The embrittlement resistance of low carbon steel may have been due to its thermalmechanical treatment history. It was the only metal of these four irons that mas received and tested as hot-rolled sheet.
I N D U S T R I A L A N D E N G I N E E R I N G CHEMISTRY
Vol. 48, No. 5
EXTREME CONDITION PROCESSING I
In
0 2
---1 ---
20'
0
I 0 U
U
0
\
15 UJ d
m
5
KEY:
z 1 0 _SS430-LOT I
I
8 7,000 PSI
E 2
6 15.000 PSI d 36,000 PSI -0 45,000 PSI d52,W PSI 0 60,OqO PSI
I
\ 16
- SQUARE
25
36
49
4
I
9
TIME IN OAYS
ROOT SCALE
Figure 3. Ductility of stainless steel Type 430 exposed to hydrogen gas a t room temperature
KEY: SS 420 LOT 2 c3 7.500 PSI 0 46.000 PSI @ 15.000 PSI 652,000 P S I d36,OOO PSI 60,000 PSI 16 940.0,OO 25 PSI 36 49 I 64 I
-
OO 9
TIME IN DAYS
5m
10
Figure 4.
,
- SQUARE
ROOT SCALE
Ductility of stainless steel Type 420 exposed to hydrogen gas at room temperature
That this effect may be a t least in part due to surface changes in low carbon steel is suggested by the fact that this metal was embrittled under acid pickling conditions. Besides low carbon steel, only Carilloy was completely unchanged by the hydrogen treatment. All other materials were embrittled to a greater or lesser extent. Particularly severe was the effect on the titanium steel, Maxel shank steel, Puron, 1-itrenamel, and the coated steels-Yifer, Alfer, and galvanized iron. Nifer and Alfer are clad with nickel and aluminum, respectively, and are especially interesting to compare with the commercially pure metals and the high chromium group. Figure 5 shows the ductility-square root of time plot for Alfer. Its resemblance to stainless steel Type 430 is marked, but this is not surprising since it has a stable aluminum oxide outer film and a macroscopic metallic aluminum layer over the steel. For the length of time it takes the hydrogen to penetrate to the steel, no embrittlement is recorded since the aluminum laycr itself is iiot embrittled (compare with aluminum in group 3.). Conversely, nickel possesses neither an oxide film of the A1103 type, nor the resistance to embrittlemcnt that aluminum has (compare with nickel in group 1.); thus, Wfer begins to show hydrogen embrittlement very soon after exposure (Figure 6 ) Just as zinc and aluminum behaved similarly toward hydrogen, in the same way galvanized iron is like Alfer in its behavior. Puron is especially interesting because it is a very high purity iron, as shown in Table 11. I t s embrittlement is a clear indication that no large alloy addition is needed to make iron susceptible and, in particular, that with only 0.005% carbon it can be affected. The remaining metals of this group were definitely, though moderately, embrittled. AISI Types 1020, 1010, and 1095 are in this category (Type 1095 was severely embrittled in one long-time run), and Rex M-2 i@included even though its initially brittle state did not allow much change to be detectable.
statement of the solubility of a gas in a metal always requires a reference condition, usually taken as the full anneal. Since the metals used were all cold-worked to a greater or lesser extent, their dislocation densities could differ from the annealed state by a factor as great as 10,000. [Barrett ( 4 ) gives lo8 and 10'1 as the dislocation densities of annealed and cold-worked metals, respectively, ] Following Cottrell's reasoning ( 9 ) , one would expect variations of several orders of magnitude in gas absorptions measured on metals after different mechanical treatments; as the dislocation density increases, a much larger number of sites become available for interstitial positioning. In a metal with a dislocation density of 10'2 per square cm. it has been estimated (4) that about 0.1 atomic % solute would be needed t o reach an equilibrium concentration. This figure is equivalent to about 1.5 in the volume ratio units reported in Table 11.
Hydrogen Content and Embrittlement
Severely embrittled ; over 10% reduction in ductility and over 50% of maximum reduction recorded Special alloy No. 3 60,200 lb./sq. inch gas, 9.7 days 0.22 AISI 1020 15N HzSOa acid, 51 hours 0.22 Stainless steel 420 7300 lb./sq. inch gas, 40 days 0.32 Nifer 45,500 lb./sq. inch gas, 10 days 0.38 Low carbon steel 15N HaS04, 51 hours 0.62 0.67 Nichrome 52,200 lb./sq. inch gas, 40 days Nickel 58,600 lb./sq. inch gas, 20 days 0.98 Alfer 2 N HzSO4 acid, 1 daw 1.2 Galvanized iron 60,200 lb./sq. inch gas, 9.7 days 1.2 1.5 Nickel 52,200 lb./sq. inch gas, 40 days Monel 14,700 lb./sq. inch gas, 27 days 1.8 a This run was made at 70" C. ; others, whether acid or gas exposure, refer to room temperature.
In twenty particular cases, broken strip samples were submerged in mercury and heated in a simple buret apparatus patterned after the one used by Darken and Smith (11) (Figure 7 ) . I n this way the evolved hydrogen was trapped and measured. The results are presented in Table 11. Expressed as the ratio of gas volume a t standard temperature and pressure to the metal volume, the amount absorbed varied from 0.045 to 4.4with blanks of 0.00 and 0.015. The magnitude of these absorption values is approximately 100 times greater than would be expected from an extrapolation of the available low pressure-high temperature data. This difference is primarily due t o a change in reference state. As already mentioned, a Mav 1956
Table II.
Gas Collection Data Grouped According to Degree of Embrittlement
Volume Ratio Gas : Conditions of Exposure Metal Not embrittled or only slightly embrittled; less than 10% reduction in ductility and less than 30% of maximum reduction recorded Low carbon steel Control, a s received 0.015 Stainless steel 420 Control, as received 0.00 AISI 1020 2 N HzSOa acid, 6 hours, 70° C.a 0.045 AISI 1020 7300 lb./sq. inch gas, 40 days 0.055 Monel 6 N HCl acid, 44 hours 0.090 Moderately embrittled ; over 10% reduction in ductility and less than 50% of maximum reduction recorded Nifer\ 60,000 lb./sq. inch gas, 6 hours Alf er 4.4 1.2
1
INDUSTRIAL AND ENGINEERING CHEMISTRY
889
ENGINEERING, DESIGN, AND PROCESS DEVELOPMENT I
2s
1
-
-
KEY': M E R 1
7,500 PSI
@
v)
n w z m LL 0
>.
c
i u t-
n a I
I
4
9
TIME IN DAYS
figure 5.
16
- SQUARE
25
36
49
ROOT SCALE
Ductility of Alfer exposed to hydrogen gas at room temperature
The values reported are of course not true solubilities, since no attempt was made to reach equilibrium, but they do indicate approximately the quantities of gas involved in the enibrittlement effect. To our knowledge these are the only measurements ever made in the very high pressure region. Perhaps the most important question that these data could shed light on is this: How is embrittlement related to the quantity of hydrogen absorbed? Because of the great number of simultaneous variables that influence gas absorption, it n-as found that a simple quantitative answer to this question is not contained in the limited data obtained; nevertheless, several qualitative classes can be noted. The data of Table I1 show that T-iorse embiittlement and greater gas absorption go hand in hand. Without exception, gas evolved from the first group of samples was less than 0.1 on the volume ratio scale that is reported. On the other end of the scale, badly affected specimens in no case gave absorption values less than 0.22 and usually evolved considerably greater amounts. The exceptions are the surprisingly high absorption values for moderately embrittled Nifer and Blfer. Actually, these seemingly anomalous data are consistent with the previous discussion of the role of surface cladding in embrittlement. As long as the absorbed hydrogen was concentrated in the relatively insensitive cladding material (compared to the underlying steel), the composite metal was only moderately embrittlecl, but after a longer time of exposure enough hydrogen entered the base steel from the IOY', cladding layer on each side to cause severe emhrittlement of the composite. Results of Bend Tests. As has already been suggested, the most important single variable controlling room temperature enibrittlemeiit is surface condition. Apparently, oxide surface films can temporarily keep hydrogen from penetrating into metals and prevent loss of dnctilitv; nevertheless, for long-time exposures Euch materials are not satisfactory, because Then the surface is penetrated by hydrogen, as it is under high pressures, the effects may be very severe. Assuming that there is no m y to permanently poison a metal surface to hydrogen, a more dependable method of design m-odd be to choose a material which retains sufficient ductility even when saturated with hydrogen. This condition n ould apply for materials like stainless steel Types 420 and 321 I n follomhg this basis for design it should be noted that the results agree n i t h previous work on embrittlement-namely, that annealed and hot-rolled metals are more satisfactory than those in the coldworked condition. This dependence on thermal and mechanical treatment may be due either to the relative porosity of the sur-
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TIME
Figure 6.
IN
DAYS
-
SQUARE ROOT SCALE
Ductility of Nifer exposed to hydrogen gas at room temperature
face film or to the bulk-metal condition where cold tvork changes the number and character of dislocations. Much of the preceding material has supported the view that the rate of embrittlement is controlled by surface reactions, hut the fact that this is true in many cases should not lead to the conclusion that it is universally true. In the argument that follows, reasons are presented for believing that a t least for some metals the surface reactions are indeed negligible and the enibrittlement rate is controlled by the diffusion step. Evidence for a connection between the equations of classical diffusion theory and the data of Figure 2 on blonel can be de-
Figure 7. A. B.
C. D. E.
F.
Gas collection apparatus
Inverted g r a d o a f e d buret Bent screen (to k e e p specimens from rising in mercury) Mercury in glass b e a k e r Bunsen burner Thermometer M e t a l specimen
INDUSTRIAL AND ENGINEERING CHEMISTRY
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EXTREME CONDITION PROCESSING veloped by replotting the Monel data on new semilog coordinates of time versus a log function of fractional ductility loss (Figure 8). The slope of the resultant straight line can be used t o calculate the diffusivity of hydrogen in Monel at room temperature (from Equation 8), if, and only if, the following are true: 1. The rate of permeation is a diffusion-controlled processthat is, surface reactions, if any, are rapid in comparison. 2. T h e loss of ductility as measured by the bend test is proportional to the amount of hydrogen absorbed. The fractional ductility loss must then be equal t o the fractional hydrogen saturation.
It is primarily as a test for these hypothese: that the mathematical treatment was attempted. Assumption 1 is necessary to validate the use of diffusion theory. Obviously the equations would be totally inadequate if any other terms in t h e over-all permeation process were significant. Experimental justification of the second assumption would be a worth-while project in itself. As it stands, however, its use requires only that a known ductility correspond to the hydrogen-saturated state. This can be read from a graph such as Figure 2. Then without measuring any hydrogen contents at all, the fractional approach t o the known saturation ductility, calculated from the bend test data, is taken as equal to F = - Ci)/(C* - Ci) of Equation 6. Though these two are fairly sweeping assumptions, they are not unreasonable; furthermore, good agreement in the literature with the diffusivity so calculated is strong evidence t h a t they are valid. Since, as explained earlier, ductility indexes frequently disagree, it is valuable t o be able to relate the bend test to a more fundamental quantity. Brief examination of Equation 8 will show that if log (1 - F ) is plotted against time, t, the curve should straighten out a t large t values, giving a line of slope proportional to the diffusivity, D. This is illustrated in Figure 8. The calculated value for hlonel is 1.2 X sq. cm./sec. This may be sq. cm./sec. given by compared with a value of D = 0.86 X Barrer (8) for hydrogen in nickel at 25” C. Other metals besides Monel gave data consistent with an explanation of negligible surface reaction. These were nickel, Xichrome V, and 35% cold-reduced I-Iaynes alloy 25; however, in none of these cases were the measurements as extensive as for RIonel, and it is believed that no conclusions should be drawn for them on this basis.
(c
Table 111.
Test
Typical Tensile Test Results
Hydrogen Gas Tensile Exposure Conditions Strength, Atm. C. Weeks Lb./Sq.Inch
Ea,
Rb,
%
%
Annealed Hypernik: Nominal Composition (50% Ni, 0.90% Mn, 0.50% Co, 0.45% Si, 0.60% all Others, Balance Fe) 25 5 78,800 331/2 63 L-0 1000 L-1 1000 25 5 79,200 301/2 61 L-2 1000 400 1 75,000 7’/e 33 L-3 2000 200 3 80,200 32’/2 64 L-4 1000 400 1 80,700 38 L-5 2000 400 1 59,800 6 29 L-6 Control 400 1 35,900 28l/2 60 L-7 2000 400 1 57,100 5 29 Special Alloy 4, Melted and Cast in Vacuum, Forged and Rolled at 1050’ C. : Nominal Composition (0.15% C, 10% Cr, 90% Fe) F4-2 F4-3 F4-4 F4-5 F4-6 F4-7
1000 Control 1000 1000 2000 2000
400 400 450 400 500 500
1 1 1 1 5Days 5Days
167,500 168,000 150,500 167,000 96,000 98,400
11 111/z 9 8 15 16
36 35 43 15l/z 54 56
Elongation in 2-inch-gage length. b Reduction in area. based on actual outside diameter6 before and after testing. c Specimen necked and broke at gage mark. a
May 1956
ob
I
10
I
ao I
20
40 I
I 50
60
TIME IN DAYS
Figure 8.
Ductility of Monel exposed to hydrogen gas at room temperature
As emphasized before, our own data indicate t h a t surface reactions are usually not negligible, denying the possibility of applying Assumption 1 t o all cases. Nevertheless, for individual metals such reasoning may be wholly valid. Kinetic data are too easy to misinterpret by curve fitting and for this reason, in spite of the close agreement on diffusivities, the picture of diffusion controlling embrittlement for even a single metal must be viewed with some reserve. While rejecting the label of “conclusive proof,” the data still provide strong evidence in favor of the two very important assumptions listed.
Tensile Tests-High
Temperature Exposure
By drilling ‘/*-inch-diameter holes in standard 1/2-inch-diameter tensile specimens, miniature pressure vessels were obtained with outside to inside diameter ratio of four. T h e hydrogen gas was compressed directly into these holes and a group of six samples together with holders and manifold were placed in a n electrically heated furnace (60). After a given exposure period, the specimens were removed and broken in tension at room temperature to obtain: 1. A complete stress-strain curve for the elastic region 2. Per cent elongation in 2-inch-gage length 3. Per cent reduction in area 4. Tensile strength
Typical results for the last three items are shown in Table I11 and one stress-strain curve is shown in Figure 9. When the hydrogen effect was embrittlement alone, only the two ductility indexes, per cent reduction in area and per cent elongation, were seen to change. This was the result in the vast majority of specimens where some property changes were observed. The result of chemical attack, much less common in our tests, was t o reduce the tensile strength of the material and alter the stress-strain curve. Since the product of chemical change is of a new chemical composition, this is not unexpected. I n most instances where attack was present, i t was accompanied by serious embrittlement, indicating that physical stresses remained even after considerable chemical reaction had occurred, but in some cases the chemical changes were severe enough to alter the
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ENGINEERING, DESIGN, A N D PROCESS DEVELQPMENT metal to a more ductile form. As a result, there was a \Tide region of overlap, and it was not always clear whether a serious effect was embrittlement or attack or both. The techniques used to differentiate attack from embrittlement xere of tx-o kindsmetallographic microscopic examination and recovery treatment. Heating, say a t 400" C. for 24 hours, could restore embrittled metals to their ductile condition, so long as they were not attacked. I n fact, the specimens could recover spontaneously a t room temperature if one waited for a prolonged peiiod before testing (10 to 50 days).
F4-6, F4-7 -t--
-
STRESS STRAIN DIAGRAM SPECIAL ALLOY 4
3
4 STRAIN,
Figure 9.
5
6
7
8
9
INGHES/INCH x 1000
Stress-strain curves for special alloy No. 4
F4 designations refer to specimens listed in Table 111; curves for specimens F4-2 and F4-4 (between F4-5 and F4-3) are omitted for clarity
Group 1 : Low-Alloy Steels. Although the inadequacy of l o w alloy steels for hydrogen service a t high temperatures had been well established a t lower pressures when this investigation was begun, a set of 16 low-alloy steels was used t o determine the effects of several variables on materials known to be embrittled and/or attacked. These metals all had carbon contents between 0.17 and 0.46'%, Chemical decarburization, therefore, R as the mode of attack. Embrittlement could not long be an isolated factor for these materials, because a t high temperatures the fairly rapid attack could soon overshadow the relatively smaller change; yet the results do indicate that embrittlement occurred before attack became severe. In some cases it n-as possible t o partially recover a specimen's ductility by heating it in air a t 400' C. for 24 hours, an obvious impossibility had it been chemically altered in an irreversible reaction. During the short exposure times used-a matter of days-the influence of temperature vas much more pronounced than that of pressure. Frequently a change of 25' C. was sufficient t o change a result from complete resistance to hydrogen t o catastrophic failure, while doubling the pressure from 1000 to 2000 atm. only moderately increased the hydrogen effect. Small additions of chromium and vanadium t o 0.2% carbon steel could not prevent embrittlement or attack entirely but did appreciably increase the useful life of a specimen. Vanadium is especially effective in this regard: At 400" C. and 1000 atm., for example, 0.08% vanadium was sufficient to protect the steel for a week, but 0.01% vanadium steel, otherwise the same in composition, was severely embrittled and attacked after 2 days a t 375' C. and 1000 atm. The ranges of compositions tested were 0.28 t o 1.49% chromium and 0.01 to 0.15y0vanadium. For each group, the protective effect Tyas evident. 892
Group 2 : Stainless Steels and High Chromium Alloys. This group, consisting of six special alloys and five commercial grade stainless steels, parallels very closely group 4 of the room temperature embrittlement tests. The six iron-chromium special alloys are identical to those used before, coming from the same melt. The five stainless steels were from new sources but nominally also quite similar. As v a s the case for the six iron-chromium special alloys when tested a t room temperature, high chromium content did not confer embrittlement resistance merely by its presence. With chromium contents from 10 to 20%, only one of the six alloys did not show any embrittlement, but because of its lower tensile strength, the conditions of exposure for this metal were somewhat milder than for the others. Further, comparison of results for similar compositions n-ith and vithout carbon gives no evidence that carbon content of 0.15% has any affect. The other high chromium alloys are all ferritic stainless steels. The austenitic stainless steels are not represented because the previous workers on this project (50) established their excellent resistance t o high temperature hydrogen gas. Stainless Type 410 was obtained 17-ith t n o different histories. It is therefore possible to compare the centerless ground alloy with hot-rolled, annealed metal. From the difference in properties between the two control samples, it is evident that the centerless grinding referred to by the manufacturer is not merely a surface treatment. If i t were only that, the preparation of specimen§ from the bar stock would be enough t o eradicate any difference. The hydrogen exposure affected the two materials in a specific manner, a reflection of the variation in manufacturrr'e treatment. Though there \vas no sign of embrittlement for the hot-rolled, annealed metal exposed at 500" C. for 2 m e k s to 2000-atm. gas, the same alloy as centerless ground mateiial was definitely embrittled by even mildcr conditions. Increase of chromium from 10 to 20% in the special alloys or from 11.8 to 27.5% in the stainless steels did not alter the hydrogen-resistant properties to any noticeable extent, as far as pmbiittlement is concerned; the piimaiy benefit t o be derived fiom chromium seems to be protection against decarburization. Other than the behavior of T l p e 410 discussed, as a group all the straight-chromium stainless steels are resistant to hydrogen over the range of conditions tested-up to 2 weeks a t 500" C. under 2000-atm. pressure in the severest test. Though Type 329 is a ferritic stainless steel, it has over 4% nickel; its severe embrittlement in contrast t o the other stainless steels appeaxs to be attributable t o this addition. Group 3 : Copper Alloys. Because copper alloys have exrellrnt resistance t o room temperature embrittlement, several highstrength copper alloys were obtained for high temperature te They were primarily combinatiotis of copper-zinc, coppc raluminum, and copper-zinc-aluminum. All the alloys of this group were resistant to hydrogen. Unfortunately, because of the danger of failure involving copper parts weakened by high teniperature, these materials could not be expoeed t o the highest pressures and temperatures. Group 4 : Armco Iron, Nickel, Boron Steel, Hyperco, Hypernik. Because steel is decarburized by hydrogen, the exact role of carbon in relation to embrittlement is difficult to establish. rlrmco iron is a pure grade of commercial iron with a very low carbon content, around 0.0370. I n agreement vith the bend test, tbe tensile test also indicated that Armco iron is embrittled at rooin temperature by 1000-atm. hydrogen and progressive13 more a t higher temperatures. Evidently, the embrittlement is not dependent on large carbon contents in iron. Grade A nickel and Hypernik Twie both embrittled in the high temperature runs, again re-enforcing the conclusion derived fiom the bend tests, that nickel and its alloys are embiittled in high pressure hydrogen service. However, Hypernik Tvas not embrittled in the tests a t room temperature. The only plausible explanation for this seems to be the presence of a protective film
INDUSTRIAL A N D E N G I N E E R I N G C H E M I S T R Y
Vol. 48, No. 5
EXTREME CONDITION PROCESSING a t room temperature which becomes less effective as the temperature is increased. Hyperco, the cobalt-iron alloy, was too brittle, as received, t o be of any use in detecting embrittlement; however, it did resist attack by 400’ C., 1000-atm. hydrogen and showed no change after 30 hours. Two boron steels were included in these tests because of earlier indications that boron-containing steels were quite resistant t o hydrogen (60). I n these earlier tests the boron steels also contained titanium, and it was realiaed that the resistance might be due to this element. T o settle this point we tested boron steels without titanium and they exhibited no greater resistance t o hydrogen than would be expected from their chromium and molybdenum contents. It is concluded that boron does not have any appreciable effect on the resistance t o hydrogen attack. Results of Tensile Tests. Previous work ( 3 7 )has shown that the standard tensile test is not a consistent index t o hydrogen content, and it would therefore not be expected t o correlate neatly with embrittlement. With this in mind, the conditions for high temperature exposure were chosen with the aim of straddling the area of effect-that is, the goal was to find the range of conditions where “no effect’’ merged int,o embrittlement and/or attack. These results, although recorded numerically, are therefore much more qualitative than the results of the bend tests, primarily because the nature of the tensile test is such that a very large number of tests would have been necessary for statistical evaluation. The bend test and tensile test agree in results, and in the region of exposure conditions where they overlap in measurement there are no outstanding contradictions between the tests. From the point of view of service design for high temperature uses, the austenitic stainless steels still are the most promising. The other alloys t h a t performed satisfactorily with respect t o hydrogen had serious drawbacks, the copper alloys being too weak and the ferritic stainless alloys as a group unpredictable. Individual alloys of the group, notably Types 420,416, and 430, appear to be satisfactory even though 420 and 430 were clearly embrittled in the room temperature tests. In this regard several metals t h a t apparently lasted without change in high temperature service, were definitely embrittled at room temperature. These are, especially, the stainless steels, austenitic as well as ferritic, but include also Nichrome. This difference could possibly be attributed to a basic difference between the two tests used, but as pointed out they did agree wherever conditions of exposure were comparable. An explanation which seems more suitable is that the metal has changed enough in the heating process t o alter its response t o a hydrogen atmosphere. Such radical changes are not unknown and were recently reported by Gulbransen and Andrew (16)for the hydrogen-zirconium system. However they reported a reverse effect and interpreted the increase of reaction rate with temperature as due to a progressive solution of the protective surface-oxide film in the metallic phase. If we are t o explain our results similarly, it is necessary to imagine a protective film that increases with temperature. For the high chromium alloys involved this may be valid but, as yet, must remain just a speculation. One other reference deserves mention in this regard: Armbruster (1) describes a marked surface-inhibiting effect for hydrogen solution in a 13% chromium steel exposed a t atmospheric pressure and elevated temperatures. The relative importance of temperature changes in this series of runs should be emphasized. Whether the mechanisms of change involved diffusion phenomena or chemical reaction rates, they would be exponential temperature functions. On the other hand, pressure effects depend on smaller rate-of-change functions, probably square roots. The experimental evidence supports these views and indicates that the choice of an alloy for high temperature-high pressure service is dependent mainly on the temperature, providing the mechanical strength of the metal is sufficient to hold the pressure. May 1956
Acknowledgment H. R. Spendelow, Jr., provided the six special iron-chromium alloys described, and his help is gratefully acknowledged. Thanks are due also t o the following organizations for metals supplied us: United States Steel Co., International Nickel Co., DriverHarris Co., Mallory-Sharon Co., Rem-Cru Titanium, Haynes Stellite Co., Metals and Controls Corp., Inland Steel Co., Crucible Steel Co. of America, Carpenter Steel Co., and Bridgeport Brass Co.
Literature Cited Armbruster, M. H., J . Am. Chem. SOC.65, 1043 (1943). Barrer, R. M., “Diffusion in and through Solids,” Cambridge Univ. Press, London, 1941. Barrer, R. M., Trans. Faraday SOC.36, 1235 (1940). Barrett, C. S., “Structure of Metals,” McGraw-Hill. New York, 1952.
Bodenstein, M., 2. Elektrochem. 28, 517 (1922). Boudart, M., J . Am. Chem. SOC.74, 3556 (1952). Bridgman, P. W., Proc. Am. Acad. Arts Sci. 59, 173 (1924). Chem. Eng. News 30, 2942 (1952). Cottrell, A. H., “Progress I n Metal Physics,” vol. 1, Pergamon Press, London, 1949. Cox, J. L., Chem. and Met. Eng. 40, 405 (1933). Darken, L. S., Smith, R. P., Corrosion 5 (January 1949). Dowden, D. A,, J . Chem. SOC.1950, p. 242. Edwards, C., J. Iron and Steel Inst. (London) 110, 9 (1924).
Gillett, H. W., “Behavior of Engineering Metals,” Wiley, New York, 1951.
Gulbransen, E. A., Andrew, K. F., J . Electrochem. Soc. 101, 348 (1954).
Ihrig, H. K., IND. ENQ.CREM.41, 2516 (1949). Inglis, N. P., Andrews, W., J . Iron and Steel Inst. (London) 128, 383 (1933).
Jacque, L., Compt. rend. 203, 936 (1936). Jaswon, M. A., “Theory of Cohesion,” Pergamon Press, London, 1954.
Johnson, W. H., Proc. Roy. SOC.23, 168 (1875). Kaainczy, F. de, J . Iron and Steel Inst. (London) 177, 85 (1954). Kosting, P. R., Metals and Alloys 5 , 54 (1934). Langdon, S. C., Grossman, M. A., Trans. Am. Electrochem. SOC. 37, 543 (1920).
Maxwell, H. L., Trans. Am. SOC.Met. 24, 213 (1936). Morris, T., J . SOC.Chem. Ind. 54, 7 (1935). Naumann, F. K., Stahl und Eisen 57, 889 (1937). Ibid., 58, 1239 (1938). Nelson, G. A., Proc. Am. Petroleum Inst. 29M(III), 163 (1949). Nelson, G. A., Trans. Am. SOC.Mech. Engr. 73, 205 (1951). Orowan, E., 2. Physik 89, 634 (1934). Pauling, L., Proc. Roy. SOC.(London) A196, 343 (1949). Poulter, T. C., Uffelman, L., Physics 3 , 147 (1932). Roberts,, J. K., Proc. Roy. SOC.(London) A152, 445 (1935). Sarjant, R. J., Middleham, T. H., Trans. Chem. Engr. Congress, World Power Conference, London, 1, 66 (1936). Schuyten, J., Corrosion and Materials Protect. 4 , No. 5, 13 (1947).
Schwab, G. M., Trans. Faraday SOC.42, 689 (1946). Seabrook, J. B., others, Trans. Am. Inst. Mining Met. Engrs. 188, 1317 (1950).
Seitz, F., “Physics of Metals,” McGraw-Hill, New York, 1943. Selwood, P. W., “Advances in Catalysis,” vol. 3 , p. 28, Academic Press, New York, 1951. Smith, D. P., “Hydrogen in Metals,” University Press, Chicago, 1948.
Smithells, C. J., “Gases and Metals,” Wiley, New York, 1937. Taylor, G. I., Proc. Roy. SOC.(London) A145, 362 (1934). Trapnell, B. M. W., “Advances in Catalysis,” vol. 3, p. 1, Academic Press, New York, 1951. Trapnell, B. M. W., “Chemisorption,” Academic Press, New York, 1955. Vanick, J. S., Chem. and Met. Eng. 34, 489 (1927). Vanick, J. S., Proc. Am. SOC.Testing Materials 24, Part 2, 348 (1924).
Vanick, J. S., Sveshnikoff, W. W. de, Thompson, J. G., Nat’l. Bur. Standards (U.S.) Tech. Papar 361 (1927).
Vanick, J. S., Trans. Am. SOC.Steel Treating 4, 62 (1923). Ibid., 12, 169 (1927).
Van Ness, H. C., Dodge, B. F., Chem. Eng. Progr. 5 1 , 2 6 6 (1955). Wheeler, H. E., Trans. Am. Inst. Mining Met. Engr. 67, 257 (1922).
Zapffe, C. A., Trans. Am. SOC.Met. 42, 387 (1950). ACCEPTEDFebruary 24, 1956. RECEIVED for review October 13, 1955.
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