- Email: [email protected]
Highly protective films on stainless steels
Materials Science and Engineering, 42 (1980) 199 - 206
199
© Elsevier Sequoia S.A., Lausanne --Printed in the Netherlands
Highly Protective F i l m s o n S t a i n l e s s S t e e l s *
G. HULTQUIST and C. LEYGRAF
Department of Physical Chemistry, Royal Institute of Technology, S-100 44 Stockholm 70 (Sweden)
SUMMARY
A newly developed surface treatment process for chromium steels is presented. Basically it involves the annealing o f chromium-alloyed steel parts under high vacuum conditions, which produces chromium enrichment in the protective surface film with no measurable chromium-depleted zone next to the film, and it considerably increases the resistance to general or local corrosion in liquid media. To obtain optimal results the process parameters (the m o s t important o f which are the annealing temperature and the oxygen partial pressure) should be matched to the alloy composition and surface pretreatment. Particularly good protective properties are obtained when the m e t h o d is combined with a subsequent treatment, such as pickling, aimed at reducing the a m o u n t o f non-metallic sulphide inclusions at the free surface. The presentation sets out the basic principles and gives some examples o f the performance o f surface-treated F e - C r alloys and commercial stainless steels.
1. INTRODUCTION
Since the corrosion properties of a steel are determined by the condition of the steel surface, it is of interest to enrich this surface with alloying components such as chromium, nickel, titanium, aluminium or silicon which are known to enhance the corrosion resistance of the steel. The present work presents a processing technique for forming surface layers with increased alloy content on chromium-alloyed steels. It is based on the selective oxidation of chromium during low temperature annealing *Presented at the International Chalmers Symposium on Surface Problems in Materials Science and Technology, GSteborg, Sweden, June 11 - 13, 1979.
of the finished steel part in a vacuum with a certain residual pressure of oxygen. During this treatment the very thin oxide film covering the steel transforms into a more chromiumrich but still very thin (< 10 nm) surface film which exhibits good corrosion resistance to normal tap water and sea-water. The concept of selective oxidation at reduced oxygen partial pressures has been extensively studied [ 1 - 4] and has been used as a process for reducing high temperature oxidation [5, 6]. In work on chromiumalloyed steels it is reported t h a t the protective abilities originate from a more or less pure Cr20 3 layer with thickness in the micron range [2, 4] and a chromium-depleted zone in the metal next to the oxide layer. Such surface layers, mostly formed at temperatures above 800 K, can exhibit good protective ability against various types of high temperature oxidation but appear to be of no value for corrosion resistance in, for example, chloride solutions. Our experience has shown that in order to attain a good resistance to the localized corrosion of stainless steels in chloride-containing media an upper limit of the protective film thickness formed during annealing in various high vacuum conditions is of the order of 100 nm. Thus the process parameters used in our work were restricted to lower temperatures and oxygen partial pressures than those referred to above. The improved corrosion properties reported in this work are the result of a systematic variation of the high vacuum annealing parameters. On the basis of approximately 25 different combinations of annealing temperature, annealing time, oxygen partial pressure and total pressure it was possible to arrive at the following conclusion. For any chromiumalloyed steel and surface pretreatment there exists an optimum set of process parameters which m o d i f y the structure and composition
200 o f the " n a t u r a l " (air formed) surface layer of the steel into a more protective film that exhibits a max im um corrosion resistance. We describe the basic principles of the technique and give some examples of the results achieved in improving the resistance to general or localized corrosion. In what follows, the words "surface m e t h o d " or "surface treatm e n t " mean the high vacuum annealing of a chromium-alloyed steel part under optimal process conditions.
2. B A S I C P R I N C I P L E S
During the heating below 800 K of an F e - C r alloy or chromium-alloyed steel in vacuum and under an oxygen-free atmosphere a change occurs in the chemical composition o f the pre-existing oxide film. In the surface layer of the alloy a reduction of oxidized iron is observed and simultaneously an enrichment o f chromium occurs since c hr om i um has a stronger t e n d e n c y to form oxide com pounds than iron has [7 - 9]. The chr om i um surface enrichment is further p r o m o t e d if a low partial pressure of oxygen is introduced into the vacuum [ 10]. In earlier work [11] we have considered the initial oxide growth of F e - C r alloys on the basis o f simple kinetics taking into account the volume diffusion of iron and chromium in the alloy and also the supply of oxygen at various partial pressures. It was f ound that, for example, at 800 K a critical oxygen partial pressure o f ab o ut 10 -5 Pa (N m -2) exists below which the early oxide formed mainly consists of chromium oxide and above which the oxide formed mainly consists of iron oxide. For various F e - C r monocrystals it was experimentally verified by means of low energy electron diffraction and Auger electron spectroscopy that the diffusion of iron at 800 K from the alloy through the initially f or m ed chromium-rich oxide layer to the gas-oxide interface is p r o m o t e d by increasing the oxygen partial pressure in the range 10 -7 - 10 -4 Pa. If a surface film with good corrosion resistance is desired, it is of importance to avoid a surface enrichment of iron. It is therefore necessary to anneal the chromium-alloyed steel in a dilute atmosphere with an oxygen partial pressure that is lower than the critical value for the formation of iron-rich oxide for
that particular annealing temperature and alloy composition. This is illustrated in Fig. 1 by composition profiles obtained by the continuous monitoring of surface composition by Auger electron spectroscopy during sputtering with argon ions. The results are for an F e - 1 0 C r alloy and are presented in terms of the Cr/(Cr + Fe) ratio as a function of sputtering time. The variation of the oxygen Auger peak with time is n o t indicated but oxygen is always present as long as a chromium surface enrichment is found. Curve a, Fig. 1, shows the chromium profile of a mechanically polished F e - 1 0 C r alloy, indicating a slight chrom i um enri chm ent in the film prior to surface treatment. Curve b shows the chromium profile of the same alloy after a successful surface t r e a t m e n t resulting in an increase in corrosion resistance. The surface t r e a t m e n t was characterized by a balance between annealing t em perat ure and oxygen partial pressure and results in a marked chromium surface enri chm ent with no measurable chromiumdepleted zone beneath the oxide film. Curve c shows a high chrom i um surface c o n t e n t after anot her surface treatment. This time, however, no increase in corrosion resistance was obtained because the oxide film modification was t oo fast and resulted in a chromiumdepleted zone n e x t to the oxide film because o f the limited supply of chrom i um from the interior of the alloy to its surface region. Hence we conclude t hat it is the chromium c o n t e n t n o t only t h r o u g h o u t the surface film but also in a narrow region next to the film
Cr/cr+Fe(%)
Cr/cr+Fe(°/o)
I 4O
20- ~
~ !
I
.-
I
I
I
d
L
Sputtering depth (arbitrary units)
Fig. 1. Chromium profiles of an F e - 1 0 C r alloy after various surface treatments.
201
t h at is of importance for the corrosion resistance. Finally, curve d is a not he r example of an unsuccessful surface treatment. The oxygen partial pressure was t oo high for that particular annealing t e m p e r a t u r e and this led to an iron e n r i c h m e n t in f r o n t of the chromium-rich zone o f the film as indicated. This t ype of iron enr i c h m e n t was found to be detrimental to the corrosion resistance of the film. Some general remarks may be made concerning the two m os t i m p o r t a n t process parameters. The tem per a t ur e interval during annealing has a lower limit at the t e m per a t u re at which the diffusion of chr om i um from the substrate to the surface becomes t oo slow. The u p p er limit is determined by the precipitation o f ch r o mi um carbides at grain boundaries and in some cases the occurrence of phase transformations. In practice this means t h a t the annealing t e m per a t ur e for chromiumalloyed steel is between 550 and 750 K. Depending on alloy composition and surface p r e t r e a t m e n t the oxygen partial pressure should range from 10 -s to 10 ~ Pa. Such conditions can be obtained in commercially available high vacuum furnaces. In summary, for each chromium-alloyed steel with a specific surface p r e t r e a t m e n t there are certain o p t i m u m combinations of process parameters which increase its corrosion resistance significantly. The favourable result of such surface treatments is attributed to the observed surface chr om i um e n r i c h m e n t with no simultaneous c hr om i um depletion n e x t to the surface film. We present a few examples of applications o f this surface t r e a t m e n t to F e - C r alloys and chromium-alloyed steels.
3. RESULTS AND DISCUSSION 3.1. Enhanced resistance to general corrosion A low alloyed steel exposed to normal tap water suffers general corrosion attack. With increased c h r o m i u m alloy c o n t e n t there is an enhanced resistance to general corrosion. T o illustrate this, various mechanically polished (800 mesh) F e - C r alloys were exposed on a surface o f 20 mm 2 to tap water at r o o m temperature. Each sample was c onne c t e d to an inert platinum electrode which was also immersed in the tap water. The variation with
time of the current between the F e - C r alloy and the platinum electrode, which is a measure of the rate of dissolution of soluble species into the water, is shown in Fig. 2(a). Two basically different behaviours are seen. The 5 and 10% Cr alloys very soon reach constant currents which correspond to general corrosion attack. The currents of the 17 and 28% Cr alloys, however, are seen to decrease with time; this cont i nued t h r o u g h o u t the experim e n t which lasted for at least 1000 h. The last two alloys remain in their passive state because of the protective ability of the surface film which hinders any kind of corrosion attack. Upon modifying the pre-existing surface films by means of a high vacuum anneal for 2 h at 700 K with an oxygen partial pressure of about 10 -6 Pa the following results were obtained (Fig. 2(b)). In the initial stage of dissolution the currents of all F e - C r alloys were at least one order of magnitude lower than the corresponding currents w i t h o u t surface treatment. With prolonged dissolution time the 10, 17 and 28% Cr alloys remained in their passive state whereas the 5% Cr alloy reverted to the active state within 15 min of exposure. Obviously this surface t r e a t m e n t results in a sufficient chrom i um surface enri chm ent for the 10% Cr alloy (see curve b in Fig. 1 for the chrom i um profile) to reach a passive state in tap water. For the 5% Cr alloy the surface t r e a t m e n t was insufficient. Figure 2 shows the variation of the dissolution current with time for the first h o u r of exposure. The same behaviour was observed for prolonged exposure up to at least 1000 h. The logarithm of the current decreased linearly with the logarithm of exposure time in all cases where such a linear decrease was observed during the first hour of exposure. Hence the chrom i um alloy c o n t e n t of around 12% that is necessary to reach a passive state in tap water for at least 1000 h of exposure is reduced to between 5 and 10% if a successful surface t r e a t m e n t is applied. The principle of matching a convenient low pressure atmosphere to an annealing temperature suitable for the alloy composition in order to increase the c h r o m i u m film composition is further illustrated in the n e x t example, which concerns the oxygen partial pressure for annealing subsequent to a hardening t reat m ent . Curve a, Fig. 3, shows a chrom i um
202 pA/cm=
~A/¢m=
5
10 ~0-
10"
28
0.1-
0.1-
0
I
I
/
0.Or
0.1
1
(a)
'
hours
901
0.1
1
hours
(b)
Fig. 2. Dissolution current vs. time for Fe-5Cr, Fe-10Cr, Fe-17Cr and Fe-28Cr exposed to tap water at room temperature (a) without surface treatment and (b) with suitable surface treatment. u m c o n t e n t o f the film a f t e r the l a t t e r annealing results in an i m p r o v e d c o r r o s i o n resistance, as seen in Fig. 4. We c o n s i d e r t h a t these f e w e x a m p l e s have d e m o n s t r a t e d t h a t t h e resistance o f F e - C r alloys t o general c o r r o s i o n is increased a f t e r surface t r e a t m e n t .
Cr'3r+Fe (%}
80"
60"
.oil I
I
I
I
3.2. E n h a n c e d resistance to localized corrosion
T h e c o m m o n e s t m o d e o f failure o f stainless steels e x p o s e d t o n e u t r a l sea-water is b y local c o r r o s i o n a t t a c k . Such a t t a c k is f r e q u e n I
I
Sputtering depth (arbitrary units) Fig. 3. Chromium profiles of a hardened and annealed Fe-13Cr ahoy: curve a, normal annealing; curve b, modified annealing.
profile o f a m a r t e n s i t i c F e - 1 3 C r a l l o y w h i c h h a d b e e n h a r d e n e d and a f t e r w a r d s c o n v e n t i o n a i l y a n n e a l e d at 825 K w i t h o u t sufficient a t t e n t i o n being p a i d t o the a c t u a l o x y g e n partial pressure. Curve b shows a n o t h e r chrom i u m profile w h e r e the high v a c u u m a t m o s p h e r e has b e e n a d a p t e d t o the a n n e a l i n g t e m p e r a t u r e b y i n t r o d u c i n g o x y g e n t o a partial pressure o f a b o u t 10 -e Pa. T h e higher c h r o m i -
(a)
(b)
Fig. 4. Comparison of the corrosion resistance of hardened Fe-13Cr screws with (a) a normal anneal and (b) a modified anneal. Corrosion tests were carried out in acid NaCI mist (ASTM B 287) for 200 h.
203 tly observed in connect i on with crevices f o r m e d by marine growth or o t h e r deposits. In order to compare the resistance to the initiation of crevice corrosion of a variety of chromium-alloyed steels use was made of an electrolytic cell which allowed us to treat the samples potentiostatically in 0.5 M NaC1 solutions at pH 8. An artificial but easily reproducible crevice g e o m e t r y was obtained by pressing a r ubber ring against the flat mechanically polished steel sample (for f ur t her details see ref. 12). A classification of the stainless steels could be obtained by comparing their critical potential for crevice corrosion, which was d ete r m i ne d as follows. After immersion o f each sample for 1 min in the chloride solution the polarization potential was m o m e n t a r i l y switched on every 2 min and increased in steps of 100 mV each time. The current was recorded as a function of time at each potential. The critical potential for crevice corrosion was det er m i ned as the lowest potential where the current was f oun d t o increase within 2 min. In Table 1 the critical potential for crevice corrosion is shown for a n u m b e r of commercially available stainless steels. Results are given for steel surfaces after wet mechanical polishing only and after a subsequent high vacuum anneal for 2 h at 650 K with an oxygen partial pressure of about 10 -6 Pa. A comparison between the mechanically polished samples (800 mesh} with and w i t h o u t a suitable surface t r e a t m e n t shows that for every stainless steel there is an i m p r o v e m e n t in resistance to the initiation of crevice corrosion. Profile studies of surface-treated samples o f the austenitic steels show that this positive ef f ect can be at t r i but e d to a marked c h r o m i u m enrichment, a slight nickel enrich-
m e n t in the oxide films and no measurable chromium-depleted zone. For SIS 2326, titanium also is f o u n d to be enriched in the surface film, whereas m o l y b d e n u m is always depleted in the protective film on molybdenum-containing steel samples before and after surface t reat m ent . However, u p o n systematically varying the process parameters it was f o u n d that the critical potential for crevice corrosion of SIS 2343 could n o t be increased above 300 mV. The reason for this is the presence of certain non-metallic sulphide inclusions which can act as starting points for the localized corrosion attack [13, 14] and which are n o t eliminated during the surface treatment. In an a t t e m p t t o reduce such inclusions and o t h e r large surface inhomogeneities the SIS 2343 sample was immersed in a pickling solution containing 15% HNO 3 and 5% HF for 10 min. The results with and w i t h o u t surface treatm e n t after pickling and ot her pret reat m ent s are shown in Table 2. Pickling the SIS 2343 steel obviously has a more favourable influence on the resistance to initiation of crevice corrosion than annealing in a dilute atmosphere has. This is because pickling n o t only reduces the a m o u n t of sulphide inclusions but also to some e x t e n t increases the c h r o m i u m c o n t e n t o f the protective film [ 1 5 ] . The main advantage, however, is the excellent corrosion properties t hat are obtained when pickling is com bi ned with a suitable surface treatment. In this c o n n e c t i o n it should be n o t e d t hat in Table 1 the less alloyed SIS 2326 steel exhibits a corrosion resistance after surface t r e a t m e n t that is superior t o t hat of ot her more highly alloyed and surface-treated stainless steels. It is suggested that the reason for this lies in the smaller sulphide inclusion
TABLE 1 Critical potential for crevice corrosion in 0.5 M NaCI of pH 8 at room temperature Composition
SIS
Fe-13Cr Fe-17Cr Fe-25Cr-5Ni-l.5Mo Fe-18Cr-2Mo-Ti Fe-18Cr-9Ni Fe-18Cr-11Ni-2.7Mo
2302 2320 2324 2326 2333 2343
Fe-20Cr-25Ni-4.4Mo-Cu
2562
ASTM
410 430 329 304 316
After mechanical polishing (800 mesh)
After surface
(mV/SCE)
(mV/SCE)
--200 +- 25 --125 +50 +75 --25 +50
--50 ± 50 +50 +250 +600 +250 +300 ~I000
+150
treatment
204 TABLE 2
TABLE 3
Critical potential for crevice corrosion in 0.5 M NaC1 of pH 8 at room temperature with various combinations of pretreatments for SIS 2343
Induction period for the initiation of crevice corrosion at different polarizing potentials
Pretreatment
Without surface treatment
• (mV/SCE) Mechanical polishing, 800 mesh Mechanical polishing, + pickling Bright annealing
+50 ± 25 +700 ± 50 +500 ± 50
With additional surface treatment
Polarizing potential
(mV/SCE)
(mV/SCE) +300 ± 50 />1000 +750 -+ 50
c o n t e n t o f this high p u r i t y ( e x t r a low interstitial) alloy. F r o m this it follows t h a t f o r this steel the resistance t o crevice c o r r o s i o n is m o r e exclusively d e t e r m i n e d b y the q u a l i t y o f t h e p r o t e c t i v e film t h a n in the case o f SIS 2 3 4 3 . Finally, T a b l e 2 also shows t h a t the resist a n c e o f a b r i g h t - a n n e a l e d SIS 2 3 4 3 surface is f u r t h e r increased w h e n an annealing t e m p e r a t u r e o f 700 K and an o x y g e n partial pressure o f 10 -5 Pa are applied during high v a c u u m annealing. So far, all e x a m p l e s o f i m p r o v e m e n t in c o r r o s i o n resistance a f t e r surface t r e a t m e n t s have b e e n based on c o r r o s i o n tests e x t e n d i n g o v e r a p e r i o d n o t longer t h a n 1 h. T a b l e 3 illustrates t h a t the i m p r o v e m e n t in c o r r o s i o n resistance is m a i n t a i n e d in a c c e l e r a t e d t e s t s lasting f o r longer t h a n 1 h. This t a b l e s h o w s h o w t h e m e a s u r e d i n d u c t i o n p e r i o d f o r the initiation o f crevice c o r r o s i o n varies w i t h polarizing p o t e n t i a l f o r SIS 2 3 2 6 w i t h a n d w i t h o u t a s u b s e q u e n t surface t r e a t m e n t . In t h e s e e x p e r i m e n t s the SIS 2 3 2 6 steel was i m m e r s e d in 0.5 M NaC1 f o r 1 m i n , a f t e r w h i c h t i m e t h e p o t e n t i a l was m o m e n t a r i l y s w i t c h e d o n at various c o n s t a n t p o t e n t i a l s . A t each p o t e n t i a l an i n d u c t i o n p e r i o d c o u l d be established d u r i n g w h i c h n o increase in the dissolut i o n c u r r e n t as a f u n c t i o n o f t i m e was rec o r d e d . F r o m T a b l e 3 it can be inferred t h a t f o r s u r f a c e - t r e a t e d SIS 2 3 2 6 the s a m e induct i o n p e r i o d is o b t a i n e d at polarizing p o t e n t i a l s t h a t are 3 0 0 - 400 m V higher t h a n f o r unt r e a t e d surfaces. Similar results w e r e o b t a i n e d f o r t h e o t h e r steels investigated. T o s u m m a r i z e , t h e resistance o f a v a r i e t y o f c o m m e r c i a l l y available stainless steels t o t h e initiation o f crevice c o r r o s i o n in n e u t r a l NaC1 solutions is increased w h e n a surface t r e a t m e n t w i t h suitable process p a r a m e t e r s
50 100 200 300 400 500 600 700 800
SIS 2326 after mechanical polishing (800 mesh)
>24 h 3 - 6 rain 1 - 2 min 10 - 30 s 1 -10s
SIS 2326 after subsequent surface treatment
>24h 5 - 10 min 1 - 2 min 1 -10s
is p e r f o r m e d . In cases w h e r e n o n - m e t a l l i c inclusions p l a y a d e s t r u c t i v e role, e x c e l l e n t c o r r o s i o n p r o p e r t i e s can be achieved w h e n t h e surface is p i c k l e d and t h e n surface t r e a t e d . 3.3.
General r e m a r k s
O n e a d v a n t a g e o f h e a t i n g the finished c h r o m i u m - a l l o y e d steel p a r t in a high v a c u u m a t m o s p h e r e is t h a t the p r o t e c t i v e film f o r m e d is t h e s a m e o v e r t h e w h o l e surface o f t h e article, i.e. even in p r e v i o u s l y p i t t e d areas and areas w h e r e o t h e r k i n d s o f surface t r e a t m e n t based on e x t e r n a l a p p l i c a t i o n w o u l d be rend e r e d difficult b y shading effects. As has b e e n p o i n t e d o u t earlier, the substantial i m p r o v e m e n t in c o r r o s i o n p r o p e r t i e s is a t t r i b u t e d t o an e n r i c h m e n t o f c h r o m i u m in the p r o t e c t i v e film. I t is possible, h o w e v e r , t o allow o t h e r alloying c o m p o n e n t s in chrom i u m - a l l o y e d steels t o be e n r i c h e d in the p r o tective film. This is d o n e b y adjusting the o x y g e n partial pressure t o t h e annealing t e m p e r a t u r e and t o the r e a c t i v i t y o f t h a t e l e m e n t to the formation of oxide compounds. The higher t h e r e a c t i v i t y t o o x y g e n o f a given alloying e l e m e n t , t h e l o w e r should be the o x y g e n partial pressures used. F o r instance, nickel e n r i c h m e n t has b e e n achieved o n an SIS 2 3 4 3 steel [ 1 6 ] . A n o t h e r a d v a n t a g e o f this surface processing t e c h n i q u e is t h e a b s e n c e o f a n y k i n d o f p o l l u t i o n p r o b l e m . This is in c o n t r a s t to m o s t c h e m i c a l surface t r e a t m e n t s such as c h r o m i u m plating, pickling or passivation. O w i n g t o t h e r a t h e r small t h i c k n e s s o f the p r o t e c t i v e films f o r m e d (5 - 10 n m ) , the films are v u l n e r a b l e u n d e r h e a v y m e c h a n i c a l a t t a c k
205 such as scraping with a knife edge. Milder types o f friction, wear or scraping, however, do n o t seem t o h a v e a negative effect. It should be borne in mind th at o t h e r kinds of surface treatm e n t o f chromium-alloyed steel such as pickling or passivation will result in protective films o f thickness even less than 5 - 10 nm and hence also vulnerable to heavy mechanical attack. To examine the reduction in corrosion resistance u p o n heavier mechanical attack the following ex p er i m ent was carried out. On bright-annealed SIS 2343 surfaces with and w i t h o u t a subsequent surface t r e a t m ent , a cross was scratched with a needle to a depth o f a few tenths of a millimetre before the samples were exposed in neutral 0.5 M NaC1 solution. On u n t r e a t e d surfaces the critical potential f o r crevice corrosion was shown to decrease f r o m 500 + 50 mV before scratching t o 350 + 50 mV after scratching. The corresponding results for the surface-treated sample were 750 and 500 mV respectively. Local corrosion attacks were always observed to initiate at points of intersection between the inscribed lines and the area under the r u bb er ring. Hence heavy mechanical attack such as scratching of the kind described above will result in a local loss of i m p r o v e m e n t of corrosion resistance. Because of the protective and m a y b e also the healing abilities of the protective film after surface t r e a t m e n t the critical potential on surface-treated and subseq u en tly scratched surfaces will still be considerably higher than t hat on unt r e a t e d and scratched surfaces since the i m p r o v e m e n t in critical potential caused by the surface treatm e n t is still 150 mV after scratching. Additional annealing in air of surfacetreated samples has shown that the beneficial e f f e c t o f the surface t r e a t m e n t is maintained t o temperatures a p p r o x i m a t e l y 50 °C below t h at at which the surface t r e a t m e n t was performed. Hence welding in air or under an inert gas will destroy the protective ability of the surface layer. As a consequence of this and also because of the vulnerability unde r heavy mechanical attack it is suggested t hat the surface t r e a t m e n t should be the last step in the manufacturing process. 4. CONCLUSIONS Upon annealing chromium-alloyed steel parts u n d er p r o p e r conditions in high vacuum
in the t e m p e r a t u r e range 550 - 750 K and in the oxygen partial pressure range 10 -8 - 10 -5 Pa a marked increase in corrosion resistance can be achieved. This effect is always connected with an observed e n r i c h m e n t of chromium in the protective film and no noticeable chrom i um depletion in the alloy region next to the film. The examples given show t hat there is an increase in the resistance of chromiumalloyed steel samples t o general as well as to local corrosion attack. G o o d protective properties are obtained when this surface treatm e n t is combined with o t h e r surface pret reat m ent s such as pickling. The layer f o r m e d after surface t r e a t m e n t is from 5 to 10 nm thick, is characterized by a gradual change o f the chemical composition from t hat on the surface to the chemical composition in the bulk alloy and has a composition which is the s a m e over the entire surface of the steel part. It is suggested that the processing technique should serve as the last step in certain manufacturing processes of chromium-alloyed steels.
ACKNOWLEDGMENTS During the course of this study the help of Swedish industries was enlisted for advice, additional tests or the provision of test materials. Thanks are due t o N. O. BjSrk (AlfaLaval AB), B. Ros6n (Bulten-Kanthal AB), W. L e h m a n n (AB Fintlings), S. O. Bernhardsson (Sandvik AB) and J. Degerb~ck (Fagersta AB). In addition, we are grateful to the Swedish Board of Technical Development for partial financial support.
REFERENCES 1 L. E. Price and G. J. Thomas, J. Inst. Met., 63 (1938) 21. 2 J. Moreau and J. B~nard, J. Inst. Met., 83 (1954 55) 87. 3 R.P. Abendroth, Trans. Metall. Soc. AIME, 230 (1964) 1735. 4 G. C. Wood and D. P. Whittle, Corros. Sci., 7 (1967) 763. 5 U.S. Patent 2,269,601 (1942), to R. Perrin. 6 A. T. Cocks, J. F. Mathews and P. L. Surman, 6th Europ. Congr. on Corrosion, London, 1977,
p. 113.
206 7 C. Leygraf, G. SchSn and S. Ekelund, Scand. J. Metall., 2 (1973) 313. 8 R. K. Wild, Corros. Sci., 14 (1974) 575. 9 K. Asami, K. Hashimoto and S. Shimodaira, Corros. Sci., 18 (1978) 125. 10 C. Leygraf, G. Hultquist and S. Ekelund, Surf. Sci., 51 (1975) 409. 11 C. Leygraf and G. Hultquist, Surf. Sci., 61 (1976) 69.
12 G. Hultquist and C. Leygraf, Proc. 7th Int. Congr. on Metal Corrosion, Rio de Janeiro, 1978. 13 G. Wrangldn, Corros. Sci., 14 (1974) 331. 14 G. Eklund, J. Electrochem. Soc., 121 (1974)467. 15 G. Hultquist and C. Leygraf, Corrosion (Houston), in the press. 16 G. Hultquist and C. Leygraf, J. Vac. Sci. Technol., in the press.
199
© Elsevier Sequoia S.A., Lausanne --Printed in the Netherlands
Highly Protective F i l m s o n S t a i n l e s s S t e e l s *
G. HULTQUIST and C. LEYGRAF
Department of Physical Chemistry, Royal Institute of Technology, S-100 44 Stockholm 70 (Sweden)
SUMMARY
A newly developed surface treatment process for chromium steels is presented. Basically it involves the annealing o f chromium-alloyed steel parts under high vacuum conditions, which produces chromium enrichment in the protective surface film with no measurable chromium-depleted zone next to the film, and it considerably increases the resistance to general or local corrosion in liquid media. To obtain optimal results the process parameters (the m o s t important o f which are the annealing temperature and the oxygen partial pressure) should be matched to the alloy composition and surface pretreatment. Particularly good protective properties are obtained when the m e t h o d is combined with a subsequent treatment, such as pickling, aimed at reducing the a m o u n t o f non-metallic sulphide inclusions at the free surface. The presentation sets out the basic principles and gives some examples o f the performance o f surface-treated F e - C r alloys and commercial stainless steels.
1. INTRODUCTION
Since the corrosion properties of a steel are determined by the condition of the steel surface, it is of interest to enrich this surface with alloying components such as chromium, nickel, titanium, aluminium or silicon which are known to enhance the corrosion resistance of the steel. The present work presents a processing technique for forming surface layers with increased alloy content on chromium-alloyed steels. It is based on the selective oxidation of chromium during low temperature annealing *Presented at the International Chalmers Symposium on Surface Problems in Materials Science and Technology, GSteborg, Sweden, June 11 - 13, 1979.
of the finished steel part in a vacuum with a certain residual pressure of oxygen. During this treatment the very thin oxide film covering the steel transforms into a more chromiumrich but still very thin (< 10 nm) surface film which exhibits good corrosion resistance to normal tap water and sea-water. The concept of selective oxidation at reduced oxygen partial pressures has been extensively studied [ 1 - 4] and has been used as a process for reducing high temperature oxidation [5, 6]. In work on chromiumalloyed steels it is reported t h a t the protective abilities originate from a more or less pure Cr20 3 layer with thickness in the micron range [2, 4] and a chromium-depleted zone in the metal next to the oxide layer. Such surface layers, mostly formed at temperatures above 800 K, can exhibit good protective ability against various types of high temperature oxidation but appear to be of no value for corrosion resistance in, for example, chloride solutions. Our experience has shown that in order to attain a good resistance to the localized corrosion of stainless steels in chloride-containing media an upper limit of the protective film thickness formed during annealing in various high vacuum conditions is of the order of 100 nm. Thus the process parameters used in our work were restricted to lower temperatures and oxygen partial pressures than those referred to above. The improved corrosion properties reported in this work are the result of a systematic variation of the high vacuum annealing parameters. On the basis of approximately 25 different combinations of annealing temperature, annealing time, oxygen partial pressure and total pressure it was possible to arrive at the following conclusion. For any chromiumalloyed steel and surface pretreatment there exists an optimum set of process parameters which m o d i f y the structure and composition
200 o f the " n a t u r a l " (air formed) surface layer of the steel into a more protective film that exhibits a max im um corrosion resistance. We describe the basic principles of the technique and give some examples of the results achieved in improving the resistance to general or localized corrosion. In what follows, the words "surface m e t h o d " or "surface treatm e n t " mean the high vacuum annealing of a chromium-alloyed steel part under optimal process conditions.
2. B A S I C P R I N C I P L E S
During the heating below 800 K of an F e - C r alloy or chromium-alloyed steel in vacuum and under an oxygen-free atmosphere a change occurs in the chemical composition o f the pre-existing oxide film. In the surface layer of the alloy a reduction of oxidized iron is observed and simultaneously an enrichment o f chromium occurs since c hr om i um has a stronger t e n d e n c y to form oxide com pounds than iron has [7 - 9]. The chr om i um surface enrichment is further p r o m o t e d if a low partial pressure of oxygen is introduced into the vacuum [ 10]. In earlier work [11] we have considered the initial oxide growth of F e - C r alloys on the basis o f simple kinetics taking into account the volume diffusion of iron and chromium in the alloy and also the supply of oxygen at various partial pressures. It was f ound that, for example, at 800 K a critical oxygen partial pressure o f ab o ut 10 -5 Pa (N m -2) exists below which the early oxide formed mainly consists of chromium oxide and above which the oxide formed mainly consists of iron oxide. For various F e - C r monocrystals it was experimentally verified by means of low energy electron diffraction and Auger electron spectroscopy that the diffusion of iron at 800 K from the alloy through the initially f or m ed chromium-rich oxide layer to the gas-oxide interface is p r o m o t e d by increasing the oxygen partial pressure in the range 10 -7 - 10 -4 Pa. If a surface film with good corrosion resistance is desired, it is of importance to avoid a surface enrichment of iron. It is therefore necessary to anneal the chromium-alloyed steel in a dilute atmosphere with an oxygen partial pressure that is lower than the critical value for the formation of iron-rich oxide for
that particular annealing temperature and alloy composition. This is illustrated in Fig. 1 by composition profiles obtained by the continuous monitoring of surface composition by Auger electron spectroscopy during sputtering with argon ions. The results are for an F e - 1 0 C r alloy and are presented in terms of the Cr/(Cr + Fe) ratio as a function of sputtering time. The variation of the oxygen Auger peak with time is n o t indicated but oxygen is always present as long as a chromium surface enrichment is found. Curve a, Fig. 1, shows the chromium profile of a mechanically polished F e - 1 0 C r alloy, indicating a slight chrom i um enri chm ent in the film prior to surface treatment. Curve b shows the chromium profile of the same alloy after a successful surface t r e a t m e n t resulting in an increase in corrosion resistance. The surface t r e a t m e n t was characterized by a balance between annealing t em perat ure and oxygen partial pressure and results in a marked chromium surface enri chm ent with no measurable chromiumdepleted zone beneath the oxide film. Curve c shows a high chrom i um surface c o n t e n t after anot her surface treatment. This time, however, no increase in corrosion resistance was obtained because the oxide film modification was t oo fast and resulted in a chromiumdepleted zone n e x t to the oxide film because o f the limited supply of chrom i um from the interior of the alloy to its surface region. Hence we conclude t hat it is the chromium c o n t e n t n o t only t h r o u g h o u t the surface film but also in a narrow region next to the film
Cr/cr+Fe(%)
Cr/cr+Fe(°/o)
I 4O
20- ~
~ !
I
.-
I
I
I
d
L
Sputtering depth (arbitrary units)
Fig. 1. Chromium profiles of an F e - 1 0 C r alloy after various surface treatments.
201
t h at is of importance for the corrosion resistance. Finally, curve d is a not he r example of an unsuccessful surface treatment. The oxygen partial pressure was t oo high for that particular annealing t e m p e r a t u r e and this led to an iron e n r i c h m e n t in f r o n t of the chromium-rich zone o f the film as indicated. This t ype of iron enr i c h m e n t was found to be detrimental to the corrosion resistance of the film. Some general remarks may be made concerning the two m os t i m p o r t a n t process parameters. The tem per a t ur e interval during annealing has a lower limit at the t e m per a t u re at which the diffusion of chr om i um from the substrate to the surface becomes t oo slow. The u p p er limit is determined by the precipitation o f ch r o mi um carbides at grain boundaries and in some cases the occurrence of phase transformations. In practice this means t h a t the annealing t e m per a t ur e for chromiumalloyed steel is between 550 and 750 K. Depending on alloy composition and surface p r e t r e a t m e n t the oxygen partial pressure should range from 10 -s to 10 ~ Pa. Such conditions can be obtained in commercially available high vacuum furnaces. In summary, for each chromium-alloyed steel with a specific surface p r e t r e a t m e n t there are certain o p t i m u m combinations of process parameters which increase its corrosion resistance significantly. The favourable result of such surface treatments is attributed to the observed surface chr om i um e n r i c h m e n t with no simultaneous c hr om i um depletion n e x t to the surface film. We present a few examples of applications o f this surface t r e a t m e n t to F e - C r alloys and chromium-alloyed steels.
3. RESULTS AND DISCUSSION 3.1. Enhanced resistance to general corrosion A low alloyed steel exposed to normal tap water suffers general corrosion attack. With increased c h r o m i u m alloy c o n t e n t there is an enhanced resistance to general corrosion. T o illustrate this, various mechanically polished (800 mesh) F e - C r alloys were exposed on a surface o f 20 mm 2 to tap water at r o o m temperature. Each sample was c onne c t e d to an inert platinum electrode which was also immersed in the tap water. The variation with
time of the current between the F e - C r alloy and the platinum electrode, which is a measure of the rate of dissolution of soluble species into the water, is shown in Fig. 2(a). Two basically different behaviours are seen. The 5 and 10% Cr alloys very soon reach constant currents which correspond to general corrosion attack. The currents of the 17 and 28% Cr alloys, however, are seen to decrease with time; this cont i nued t h r o u g h o u t the experim e n t which lasted for at least 1000 h. The last two alloys remain in their passive state because of the protective ability of the surface film which hinders any kind of corrosion attack. Upon modifying the pre-existing surface films by means of a high vacuum anneal for 2 h at 700 K with an oxygen partial pressure of about 10 -6 Pa the following results were obtained (Fig. 2(b)). In the initial stage of dissolution the currents of all F e - C r alloys were at least one order of magnitude lower than the corresponding currents w i t h o u t surface treatment. With prolonged dissolution time the 10, 17 and 28% Cr alloys remained in their passive state whereas the 5% Cr alloy reverted to the active state within 15 min of exposure. Obviously this surface t r e a t m e n t results in a sufficient chrom i um surface enri chm ent for the 10% Cr alloy (see curve b in Fig. 1 for the chrom i um profile) to reach a passive state in tap water. For the 5% Cr alloy the surface t r e a t m e n t was insufficient. Figure 2 shows the variation of the dissolution current with time for the first h o u r of exposure. The same behaviour was observed for prolonged exposure up to at least 1000 h. The logarithm of the current decreased linearly with the logarithm of exposure time in all cases where such a linear decrease was observed during the first hour of exposure. Hence the chrom i um alloy c o n t e n t of around 12% that is necessary to reach a passive state in tap water for at least 1000 h of exposure is reduced to between 5 and 10% if a successful surface t r e a t m e n t is applied. The principle of matching a convenient low pressure atmosphere to an annealing temperature suitable for the alloy composition in order to increase the c h r o m i u m film composition is further illustrated in the n e x t example, which concerns the oxygen partial pressure for annealing subsequent to a hardening t reat m ent . Curve a, Fig. 3, shows a chrom i um
202 pA/cm=
~A/¢m=
5
10 ~0-
10"
28
0.1-
0.1-
0
I
I
/
0.Or
0.1
1
(a)
'
hours
901
0.1
1
hours
(b)
Fig. 2. Dissolution current vs. time for Fe-5Cr, Fe-10Cr, Fe-17Cr and Fe-28Cr exposed to tap water at room temperature (a) without surface treatment and (b) with suitable surface treatment. u m c o n t e n t o f the film a f t e r the l a t t e r annealing results in an i m p r o v e d c o r r o s i o n resistance, as seen in Fig. 4. We c o n s i d e r t h a t these f e w e x a m p l e s have d e m o n s t r a t e d t h a t t h e resistance o f F e - C r alloys t o general c o r r o s i o n is increased a f t e r surface t r e a t m e n t .
Cr'3r+Fe (%}
80"
60"
.oil I
I
I
I
3.2. E n h a n c e d resistance to localized corrosion
T h e c o m m o n e s t m o d e o f failure o f stainless steels e x p o s e d t o n e u t r a l sea-water is b y local c o r r o s i o n a t t a c k . Such a t t a c k is f r e q u e n I
I
Sputtering depth (arbitrary units) Fig. 3. Chromium profiles of a hardened and annealed Fe-13Cr ahoy: curve a, normal annealing; curve b, modified annealing.
profile o f a m a r t e n s i t i c F e - 1 3 C r a l l o y w h i c h h a d b e e n h a r d e n e d and a f t e r w a r d s c o n v e n t i o n a i l y a n n e a l e d at 825 K w i t h o u t sufficient a t t e n t i o n being p a i d t o the a c t u a l o x y g e n partial pressure. Curve b shows a n o t h e r chrom i u m profile w h e r e the high v a c u u m a t m o s p h e r e has b e e n a d a p t e d t o the a n n e a l i n g t e m p e r a t u r e b y i n t r o d u c i n g o x y g e n t o a partial pressure o f a b o u t 10 -e Pa. T h e higher c h r o m i -
(a)
(b)
Fig. 4. Comparison of the corrosion resistance of hardened Fe-13Cr screws with (a) a normal anneal and (b) a modified anneal. Corrosion tests were carried out in acid NaCI mist (ASTM B 287) for 200 h.
203 tly observed in connect i on with crevices f o r m e d by marine growth or o t h e r deposits. In order to compare the resistance to the initiation of crevice corrosion of a variety of chromium-alloyed steels use was made of an electrolytic cell which allowed us to treat the samples potentiostatically in 0.5 M NaC1 solutions at pH 8. An artificial but easily reproducible crevice g e o m e t r y was obtained by pressing a r ubber ring against the flat mechanically polished steel sample (for f ur t her details see ref. 12). A classification of the stainless steels could be obtained by comparing their critical potential for crevice corrosion, which was d ete r m i ne d as follows. After immersion o f each sample for 1 min in the chloride solution the polarization potential was m o m e n t a r i l y switched on every 2 min and increased in steps of 100 mV each time. The current was recorded as a function of time at each potential. The critical potential for crevice corrosion was det er m i ned as the lowest potential where the current was f oun d t o increase within 2 min. In Table 1 the critical potential for crevice corrosion is shown for a n u m b e r of commercially available stainless steels. Results are given for steel surfaces after wet mechanical polishing only and after a subsequent high vacuum anneal for 2 h at 650 K with an oxygen partial pressure of about 10 -6 Pa. A comparison between the mechanically polished samples (800 mesh} with and w i t h o u t a suitable surface t r e a t m e n t shows that for every stainless steel there is an i m p r o v e m e n t in resistance to the initiation of crevice corrosion. Profile studies of surface-treated samples o f the austenitic steels show that this positive ef f ect can be at t r i but e d to a marked c h r o m i u m enrichment, a slight nickel enrich-
m e n t in the oxide films and no measurable chromium-depleted zone. For SIS 2326, titanium also is f o u n d to be enriched in the surface film, whereas m o l y b d e n u m is always depleted in the protective film on molybdenum-containing steel samples before and after surface t reat m ent . However, u p o n systematically varying the process parameters it was f o u n d that the critical potential for crevice corrosion of SIS 2343 could n o t be increased above 300 mV. The reason for this is the presence of certain non-metallic sulphide inclusions which can act as starting points for the localized corrosion attack [13, 14] and which are n o t eliminated during the surface treatment. In an a t t e m p t t o reduce such inclusions and o t h e r large surface inhomogeneities the SIS 2343 sample was immersed in a pickling solution containing 15% HNO 3 and 5% HF for 10 min. The results with and w i t h o u t surface treatm e n t after pickling and ot her pret reat m ent s are shown in Table 2. Pickling the SIS 2343 steel obviously has a more favourable influence on the resistance to initiation of crevice corrosion than annealing in a dilute atmosphere has. This is because pickling n o t only reduces the a m o u n t of sulphide inclusions but also to some e x t e n t increases the c h r o m i u m c o n t e n t o f the protective film [ 1 5 ] . The main advantage, however, is the excellent corrosion properties t hat are obtained when pickling is com bi ned with a suitable surface treatment. In this c o n n e c t i o n it should be n o t e d t hat in Table 1 the less alloyed SIS 2326 steel exhibits a corrosion resistance after surface t r e a t m e n t that is superior t o t hat of ot her more highly alloyed and surface-treated stainless steels. It is suggested that the reason for this lies in the smaller sulphide inclusion
TABLE 1 Critical potential for crevice corrosion in 0.5 M NaCI of pH 8 at room temperature Composition
SIS
Fe-13Cr Fe-17Cr Fe-25Cr-5Ni-l.5Mo Fe-18Cr-2Mo-Ti Fe-18Cr-9Ni Fe-18Cr-11Ni-2.7Mo
2302 2320 2324 2326 2333 2343
Fe-20Cr-25Ni-4.4Mo-Cu
2562
ASTM
410 430 329 304 316
After mechanical polishing (800 mesh)
After surface
(mV/SCE)
(mV/SCE)
--200 +- 25 --125 +50 +75 --25 +50
--50 ± 50 +50 +250 +600 +250 +300 ~I000
+150
treatment
204 TABLE 2
TABLE 3
Critical potential for crevice corrosion in 0.5 M NaC1 of pH 8 at room temperature with various combinations of pretreatments for SIS 2343
Induction period for the initiation of crevice corrosion at different polarizing potentials
Pretreatment
Without surface treatment
• (mV/SCE) Mechanical polishing, 800 mesh Mechanical polishing, + pickling Bright annealing
+50 ± 25 +700 ± 50 +500 ± 50
With additional surface treatment
Polarizing potential
(mV/SCE)
(mV/SCE) +300 ± 50 />1000 +750 -+ 50
c o n t e n t o f this high p u r i t y ( e x t r a low interstitial) alloy. F r o m this it follows t h a t f o r this steel the resistance t o crevice c o r r o s i o n is m o r e exclusively d e t e r m i n e d b y the q u a l i t y o f t h e p r o t e c t i v e film t h a n in the case o f SIS 2 3 4 3 . Finally, T a b l e 2 also shows t h a t the resist a n c e o f a b r i g h t - a n n e a l e d SIS 2 3 4 3 surface is f u r t h e r increased w h e n an annealing t e m p e r a t u r e o f 700 K and an o x y g e n partial pressure o f 10 -5 Pa are applied during high v a c u u m annealing. So far, all e x a m p l e s o f i m p r o v e m e n t in c o r r o s i o n resistance a f t e r surface t r e a t m e n t s have b e e n based on c o r r o s i o n tests e x t e n d i n g o v e r a p e r i o d n o t longer t h a n 1 h. T a b l e 3 illustrates t h a t the i m p r o v e m e n t in c o r r o s i o n resistance is m a i n t a i n e d in a c c e l e r a t e d t e s t s lasting f o r longer t h a n 1 h. This t a b l e s h o w s h o w t h e m e a s u r e d i n d u c t i o n p e r i o d f o r the initiation o f crevice c o r r o s i o n varies w i t h polarizing p o t e n t i a l f o r SIS 2 3 2 6 w i t h a n d w i t h o u t a s u b s e q u e n t surface t r e a t m e n t . In t h e s e e x p e r i m e n t s the SIS 2 3 2 6 steel was i m m e r s e d in 0.5 M NaC1 f o r 1 m i n , a f t e r w h i c h t i m e t h e p o t e n t i a l was m o m e n t a r i l y s w i t c h e d o n at various c o n s t a n t p o t e n t i a l s . A t each p o t e n t i a l an i n d u c t i o n p e r i o d c o u l d be established d u r i n g w h i c h n o increase in the dissolut i o n c u r r e n t as a f u n c t i o n o f t i m e was rec o r d e d . F r o m T a b l e 3 it can be inferred t h a t f o r s u r f a c e - t r e a t e d SIS 2 3 2 6 the s a m e induct i o n p e r i o d is o b t a i n e d at polarizing p o t e n t i a l s t h a t are 3 0 0 - 400 m V higher t h a n f o r unt r e a t e d surfaces. Similar results w e r e o b t a i n e d f o r t h e o t h e r steels investigated. T o s u m m a r i z e , t h e resistance o f a v a r i e t y o f c o m m e r c i a l l y available stainless steels t o t h e initiation o f crevice c o r r o s i o n in n e u t r a l NaC1 solutions is increased w h e n a surface t r e a t m e n t w i t h suitable process p a r a m e t e r s
50 100 200 300 400 500 600 700 800
SIS 2326 after mechanical polishing (800 mesh)
>24 h 3 - 6 rain 1 - 2 min 10 - 30 s 1 -10s
SIS 2326 after subsequent surface treatment
>24h 5 - 10 min 1 - 2 min 1 -10s
is p e r f o r m e d . In cases w h e r e n o n - m e t a l l i c inclusions p l a y a d e s t r u c t i v e role, e x c e l l e n t c o r r o s i o n p r o p e r t i e s can be achieved w h e n t h e surface is p i c k l e d and t h e n surface t r e a t e d . 3.3.
General r e m a r k s
O n e a d v a n t a g e o f h e a t i n g the finished c h r o m i u m - a l l o y e d steel p a r t in a high v a c u u m a t m o s p h e r e is t h a t the p r o t e c t i v e film f o r m e d is t h e s a m e o v e r t h e w h o l e surface o f t h e article, i.e. even in p r e v i o u s l y p i t t e d areas and areas w h e r e o t h e r k i n d s o f surface t r e a t m e n t based on e x t e r n a l a p p l i c a t i o n w o u l d be rend e r e d difficult b y shading effects. As has b e e n p o i n t e d o u t earlier, the substantial i m p r o v e m e n t in c o r r o s i o n p r o p e r t i e s is a t t r i b u t e d t o an e n r i c h m e n t o f c h r o m i u m in the p r o t e c t i v e film. I t is possible, h o w e v e r , t o allow o t h e r alloying c o m p o n e n t s in chrom i u m - a l l o y e d steels t o be e n r i c h e d in the p r o tective film. This is d o n e b y adjusting the o x y g e n partial pressure t o t h e annealing t e m p e r a t u r e and t o the r e a c t i v i t y o f t h a t e l e m e n t to the formation of oxide compounds. The higher t h e r e a c t i v i t y t o o x y g e n o f a given alloying e l e m e n t , t h e l o w e r should be the o x y g e n partial pressures used. F o r instance, nickel e n r i c h m e n t has b e e n achieved o n an SIS 2 3 4 3 steel [ 1 6 ] . A n o t h e r a d v a n t a g e o f this surface processing t e c h n i q u e is t h e a b s e n c e o f a n y k i n d o f p o l l u t i o n p r o b l e m . This is in c o n t r a s t to m o s t c h e m i c a l surface t r e a t m e n t s such as c h r o m i u m plating, pickling or passivation. O w i n g t o t h e r a t h e r small t h i c k n e s s o f the p r o t e c t i v e films f o r m e d (5 - 10 n m ) , the films are v u l n e r a b l e u n d e r h e a v y m e c h a n i c a l a t t a c k
205 such as scraping with a knife edge. Milder types o f friction, wear or scraping, however, do n o t seem t o h a v e a negative effect. It should be borne in mind th at o t h e r kinds of surface treatm e n t o f chromium-alloyed steel such as pickling or passivation will result in protective films o f thickness even less than 5 - 10 nm and hence also vulnerable to heavy mechanical attack. To examine the reduction in corrosion resistance u p o n heavier mechanical attack the following ex p er i m ent was carried out. On bright-annealed SIS 2343 surfaces with and w i t h o u t a subsequent surface t r e a t m ent , a cross was scratched with a needle to a depth o f a few tenths of a millimetre before the samples were exposed in neutral 0.5 M NaC1 solution. On u n t r e a t e d surfaces the critical potential f o r crevice corrosion was shown to decrease f r o m 500 + 50 mV before scratching t o 350 + 50 mV after scratching. The corresponding results for the surface-treated sample were 750 and 500 mV respectively. Local corrosion attacks were always observed to initiate at points of intersection between the inscribed lines and the area under the r u bb er ring. Hence heavy mechanical attack such as scratching of the kind described above will result in a local loss of i m p r o v e m e n t of corrosion resistance. Because of the protective and m a y b e also the healing abilities of the protective film after surface t r e a t m e n t the critical potential on surface-treated and subseq u en tly scratched surfaces will still be considerably higher than t hat on unt r e a t e d and scratched surfaces since the i m p r o v e m e n t in critical potential caused by the surface treatm e n t is still 150 mV after scratching. Additional annealing in air of surfacetreated samples has shown that the beneficial e f f e c t o f the surface t r e a t m e n t is maintained t o temperatures a p p r o x i m a t e l y 50 °C below t h at at which the surface t r e a t m e n t was performed. Hence welding in air or under an inert gas will destroy the protective ability of the surface layer. As a consequence of this and also because of the vulnerability unde r heavy mechanical attack it is suggested t hat the surface t r e a t m e n t should be the last step in the manufacturing process. 4. CONCLUSIONS Upon annealing chromium-alloyed steel parts u n d er p r o p e r conditions in high vacuum
in the t e m p e r a t u r e range 550 - 750 K and in the oxygen partial pressure range 10 -8 - 10 -5 Pa a marked increase in corrosion resistance can be achieved. This effect is always connected with an observed e n r i c h m e n t of chromium in the protective film and no noticeable chrom i um depletion in the alloy region next to the film. The examples given show t hat there is an increase in the resistance of chromiumalloyed steel samples t o general as well as to local corrosion attack. G o o d protective properties are obtained when this surface treatm e n t is combined with o t h e r surface pret reat m ent s such as pickling. The layer f o r m e d after surface t r e a t m e n t is from 5 to 10 nm thick, is characterized by a gradual change o f the chemical composition from t hat on the surface to the chemical composition in the bulk alloy and has a composition which is the s a m e over the entire surface of the steel part. It is suggested that the processing technique should serve as the last step in certain manufacturing processes of chromium-alloyed steels.
ACKNOWLEDGMENTS During the course of this study the help of Swedish industries was enlisted for advice, additional tests or the provision of test materials. Thanks are due t o N. O. BjSrk (AlfaLaval AB), B. Ros6n (Bulten-Kanthal AB), W. L e h m a n n (AB Fintlings), S. O. Bernhardsson (Sandvik AB) and J. Degerb~ck (Fagersta AB). In addition, we are grateful to the Swedish Board of Technical Development for partial financial support.
REFERENCES 1 L. E. Price and G. J. Thomas, J. Inst. Met., 63 (1938) 21. 2 J. Moreau and J. B~nard, J. Inst. Met., 83 (1954 55) 87. 3 R.P. Abendroth, Trans. Metall. Soc. AIME, 230 (1964) 1735. 4 G. C. Wood and D. P. Whittle, Corros. Sci., 7 (1967) 763. 5 U.S. Patent 2,269,601 (1942), to R. Perrin. 6 A. T. Cocks, J. F. Mathews and P. L. Surman, 6th Europ. Congr. on Corrosion, London, 1977,
p. 113.
206 7 C. Leygraf, G. SchSn and S. Ekelund, Scand. J. Metall., 2 (1973) 313. 8 R. K. Wild, Corros. Sci., 14 (1974) 575. 9 K. Asami, K. Hashimoto and S. Shimodaira, Corros. Sci., 18 (1978) 125. 10 C. Leygraf, G. Hultquist and S. Ekelund, Surf. Sci., 51 (1975) 409. 11 C. Leygraf and G. Hultquist, Surf. Sci., 61 (1976) 69.
12 G. Hultquist and C. Leygraf, Proc. 7th Int. Congr. on Metal Corrosion, Rio de Janeiro, 1978. 13 G. Wrangldn, Corros. Sci., 14 (1974) 331. 14 G. Eklund, J. Electrochem. Soc., 121 (1974)467. 15 G. Hultquist and C. Leygraf, Corrosion (Houston), in the press. 16 G. Hultquist and C. Leygraf, J. Vac. Sci. Technol., in the press.